Nanocomposite coatings: thermal spray processing,microstructure and performance

Nanotechnology and nanomaterials

The concept of nanotechnology has emerged since 1959 when the Physics Nobel Laureate, Richard Feynman gave his renown talk ‘There's Plenty of Room at the Bottom’. ¹ However, it was not until the early 1980s that the prospects of atomic-scale control of materials became prominent through the development of a conceptual framework for nanotechnology ² and technological advances such as the invention of scanning tunnelling microscopy ³ and the discovery of fullerenes. ⁴ The research and development (R&D) in nanotechnology accelerated in 2000 when the definition and long-term vision of nanotechnology was defined by the United States National Nanotechnology Initiative. ⁵

This advancement immediately sparked the establishment of sustained nanotechnology R&D by Japan, Korea, the European Community, Germany, China and Taiwan. ⁶ Figure 1 shows the growth of the national funding on nanotechnology R&D for several countries since the mid-1990s. More detailed information and analyses on the trend of R&D funding on nanotechnology, including funding by public and private sectors across different countries and their projected growth until 2015, can be found in Refs. 6 and 7.

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Figure 1.

The federal/national government funding for research and development (R&D) in nanotechnology from 1997 to 2009. Source: World Technology Evaluation Center, Inc. ^5,6

The European Commission defined nanomaterials as ‘a natural, incidental or manufactured material containing particles, in an unbound state or as an aggregate or as an agglomerate and where, for 50% or more of the particles in the number size distribution, one or more external dimensions is in the size range 1–100 nm. In specific cases and where warranted by concerns for the environment, health, safety or competitiveness the number size distribution threshold of 50% may be replaced by a threshold between 1 and 50%.’ ⁸ In other words, nanomaterials display a structure with at least 50% (by number) of a material with at least one dimension smaller than 100 nm. Note that before the widespread use of the term nanomaterial or nanoparticle, materials in the nano-sized range; i.e. less than 1 μm, were referred to as ultrafine or sub-micrometre particles. ^9–11

Nanomaterials can be classified according to their geometric dimensions; i.e. the number of dimensions that are not confined to the nanoscale range. ¹² The four classifications of nanomaterials are (i) zero-dimensional (0D), where all three dimensions are within the nanoscale, such as nanoparticles or nanoclusters; (ii) one-dimensional (1D), where only two dimensions are within the nanoscale, such as nanotubes or nanowires; (iii) two-dimensional (2D), where only one dimension is within the nanoscale, such as thin films or nanoplates; and (iv) three-dimensional (3D), where none of the bulk dimensions are within the nanoscale but the constituents within the material have nano-characteristics, such as nanocrystalline and nanocomposite materials; Fig. 2.

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Figure 2.

Classification of nanomaterials: a zero-dimensional, b one-dimensional, c two-dimensional, and d three-dimensional nanostructures. Please note that ø and t represent the diameter and thickness, respectively

Nanomaterials demonstrate unique optical, ^13,14 thermal, ^15,16 mechanical, ^17,18 electrical, ^19,20 and magnetic ^21,22 properties that are not exhibited by their bulk counterparts. The attributes of nanostructured materials that are size-dependent arise from (i) surface effects due to a significant increase of surface-to-volume ratio and (ii) quantum effects due to confinement of electron movement, changes in the electronic structure and interatomic relation, and the presence of defects. ^12,23 Nanomaterials have found many practical applications in industry that range from electronics to aerospace, Table 1.

Generally, there are two approaches to produce nanostructured materials: (i) the top-down and (ii) the bottom-up approaches. The top-down approach involves the breakdown of a bulk material structure to reduce the crystal size to sub-micrometer or nanoscale. On the other hand, the bottom-up approach involves processing or building up a material from the atomic scale or nanoclusters in a way that retains the original nanoscale structural units. Examples of these two approaches are outlined in Fig. 3. This review concerns the deposition of nanocomposites via thermal spray route, which is a bottom-up approach. It might be argued that the impact process experienced during the formation of a thermal spray coating will cause fragmentation of micrometre-sized particles to the nanoscale; i.e. a top-down approach. However, as this is not the preferred formation process that would retain a lamellar structure, then the bottom-up approach is considered as the dominant mechanism.

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Figure 3.

The top-down and bottom-up approaches for nanomaterials fabrication

Nanocomposites

Nanocomposites are multiphase materials that are composed of two or more substances, of which at least one has a nanoscale dimension. ²⁴ The multifunctional potential of nanocomposites is manifested through combinations of the distinctive properties of each phase that act synergistically to create properties that would not otherwise be achievable. Naturally occurring nanocomposites are formed through self-organisation and directed assembly of macromolecules, inorganic materials or a combination of both in biological systems. ²⁵ Examples of natural nanocomposites include tooth enamel, bone, the aragonitic nacreous layer of abalone shell, dragline spider silk, magnetic bacteria, and sea urchin spines. ²⁵ These naturally occurring nanocomposites exhibit unique properties that are not found in synthetic materials, leading to efforts to understand the design of biological systems and create biomimetic materials such as bone scaffold and nanofibers. ²⁶

Current and potential industrial applications of nanocomposites include, among others, the field of aerospace, ²⁷ automotive, ²⁸ biomedical, ^29,30 electronics, ³¹ energy storage, ^32,33 food packaging, ³⁴ functional textiles, ³⁵ manufacturing, ^36,37 optics, ³⁸ and photocatalysis. ³⁹ The primary challenges in commercialisation of nanocomposites lies in (i) achieving precise control over the size distribution and dispersion of the nanoscale constituents for large-scale manufacturing and processing and (ii) understanding the functional roles of the interfaces between structurally and/or chemically dissimilar phases on the bulk material properties. Much research has been carried out to address these issues and explore the functionalities of nanocomposites. The development of nanocomposites can be seen worldwide, as evidenced through the increasing growth of Science Citation Index (SCI) publications and patent applications in the past 25 years; Figs. 4 and 5.

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Figure 4.

Nanocomposite publications in the Science Citation Index (SCI) from 1990 to 2012 for the five leading countries/territories. The data were generated from an online search in the Web of Science using a ‘title-abstract’ search in the SCI database for nanocomposite

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Figure 5.

Nanocomposite patent applications from 1990 to 2012 for the five leading patent offices of China (CN), United States (US), European (EU), Korea (KR), and Japan (JP). The data were generated from an online search in LexisNexis TotalPatent using a ‘title-abstract-claims’ search for nanocomposite

The upsurge of interest and research activities in the area of nanocomposites are indicated in the number of SCI publications, which reached a total of over 31 000 worldwide in 2012 compared to 1 319 in 2000. This growth is rapid and irregular around the world, with the five leading regions contributing to over 80% of the total publications; Fig. 4. The inventions related to nanocomposites are represented by the number of patent applications filed in the patent repositories worldwide. There were over 1 200 applications filed in 2012 worldwide compared to 142 in 2000. Over 75% of these patents were filed in the top five leading patent offices; Fig. 5. These statistics reflect the rapid development of nanocomposites since 2000, where the number of SCI publications and patent applications show average annual growth rates of 33 and 24%, respectively.

Despite their advantages, bulk nanocomposites can be expensive to produce because of the production costs of nanomaterials and technical limitations in preserving a nanostructure in the consolidated product due to excessive grain growth. ⁴⁰ Nanocomposite coatings use lesser amounts of nanostructured materials because they are quite thin; i.e. several micrometres up to a millimeter depending on the deposition technique, and can potentially cover large areas. Therefore, nanocomposite coatings present an effective alternative to take advantage of the many remarkable properties of nanomaterials without bearing the production costs of bulk nanocomposites. There are two types of nanocomposite coatings: (i) matrix-reinforced, where the reinforcing phase is within nanoscale, and (ii) layered coatings, where the thickness of individual layers are within nanoscale dimensions; Fig. 6.

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Figure 6.

Schematics of nanocomposite coatings: a matrix-reinforced and b multi layered

Deposition of nanocomposite coatings is a bottom up approach of producing nanomaterials. There are various techniques that can be used to manufacture nanocomposite coatings, depending on the thickness, structure, purity, crystallinity, costs, and/or materials considerations. These techniques can be classified into (i) physical methods; e.g. cathodic arc deposition, ⁴¹ electron beam evaporation, ⁴² pulsed laser deposition, ⁴³ sputter deposition, ⁴⁴ ion implantation-deposition process, ⁴⁵ and thermal spray processes; ⁴⁶ and (ii) chemical methods; e.g. chemical vapour deposition (CVD), ⁴⁷ electroplating, ⁴⁸ self-assembly, ⁴⁹ spin coating, ⁵⁰ and sol-gel dip coating. ⁵¹ Following deposition, the coatings can undergo post-deposition treatment and/or surface patterning as required for specific applications.

Thermal spray nanocomposite coatings

Thermal spray is a generic term referring to a group of coating processes in which a stream of finely divided metal or non-metal particles is melted and propelled towards a substrate to form flattened lamellae that stack progressively to form a coating. ^52,53 Thermal spray composite coatings have long been established as high performance functional coatings for applications such as wear resistance, corrosion resistance, abradable seals, and thermal barriers. ^54–57 These coatings include a variety of metal-matrix, ^58–60 ceramic-matrix, ^61,62 and polymer-matrix ^63,64 composites. Thermal spray fibre-reinforced coatings have also gained much attention since the 1980s. ^65–67 Thermal spray processes present a versatile approach for deposition of nanocomposite coatings in terms of materials, deposition medium, and thickness compared to other deposition processes.

A wide range of materials can be deposited using thermal spray processes because of its wide range of processing temperatures and velocities. Thus, any composition of materials can be combined and deposited by thermal spray to synthesise composites with optimised properties; e.g. high melting point ceramics and low melting point polymers. Thermal spray nanocomposite coatings are much thicker than those achievable by other methods; e.g. physical and chemical vapour deposited films are typically <5 and 0·1 to 500 μm thick, respectively. The thickness of thermal spray coatings range from 50 μm to several millimetres, although thinner coatings of ∼10–50 μm are possible with liquid thermal spraying.

Thermal spray processes are economically advantageous compared to other deposition methods because of (i) high deposition rate, namely <5 μm h⁻¹ for physical vapour deposition (PVD) ⁶⁸ and <200 μm h⁻¹ for plasma-assisted CVD ⁶⁹ compared to >1 000 μm h⁻¹ for plasma spray; ⁷⁰ (ii) low starting materials, equipment, and operating costs; (iii) flexibility of deposition medium; and (iv) optional substrate pre-heating. The waste disposal systems of thermal spray processes are simple and less hazardous compared to processes such as CVD and electrodeposition.

Thermal spray processes have the unique capability for dimensional restoration of damaged or worn out engineering and structural components without needing to replace the whole component. This presents a sustainable approach for incorporating nanocomposite materials as thermal barrier, wear resistant, corrosion resistant, and abradable coatings. The portability of thermal spray equipment is an advantage for on-site repair of large or immobile components. Thermal spray can, therefore, be considered as an additive manufacturing technology.

The basic structural building blocks of thermal spray coatings are splats, which are formed by the flattening of molten particles upon impact. Depending on factors such as the degree of particle melting and impact velocity, the microstructure of thermal spray coatings is typically porous, anisotropic, and may comprise microcracks, unmelted particles, and oxide inclusions. Such porous and lamellar structures are distinct from the dense columnar and/or equiaxed structures formed by atomistic deposition processes such as PVD, CVD, and ion implantation-deposition. ^69,71 Electroplated coatings exhibit grains with preferred orientation that depend on the material, electrolyte composition, deposition conditions, and substrates. ⁷² The microstructure of sol-gel coatings derived through dip coating or spin coating range from porous, weakly branched polymers to highly condensed particles, depending on the structure of the entrained inorganic species in the original sol, their reactivity, the timescale of the process, and magnitude of shear and capillary forces. ⁷³

Novel and advanced nanocomposite coatings with unique compositions and properties can potentially be deposited via thermal spray routes through the incorporation of metastable and amorphous phases formed by rapid cooling of the materials; namely 10⁵–10⁸°C s⁻¹. ⁵² Oxide-free deposits can be produced using low pressure plasma spray and cold spray, while combinations of a wide range of oxide, carbide, and nitride nanocomposites are possible through atmospheric plasma spray or through proper control of the oxygen–fuel ratio in combustion thermal spray systems. Grain growth during thermal spray processes is restricted because of the high solidification rate; thereby enabling formation of a nanocrystalline structure.

Owing to the nature of the deposition process, thermal spray coatings exhibit randomly distributed splats. Thus, there is a random phase distribution in nanocomposite systems, which is favourable in certain applications; e.g. randomly distributed (∼15%) amorphous calcium phosphate and α-tricalcium phosphate in hydroxyapatite biomedical coatings resorb preferentially to promote implant fixation. ⁷⁴ The variation of feedstock particle size distribution and morphology contributes to structural and compositional inhomogeneity in thermal spray coatings. This is especially true for nanocomposite systems because the melting and reconsolidation of feedstock materials are restricted. Therefore, the compositional homogeneity of nanocomposite coatings is reflected in the homogeneity of the original feedstock.

In terms of its limitations, the thermal spray process is a line of sight process, which infers that expensive, automation systems are necessary for deposition onto complex geometries. Thermal spray coatings are porous, exhibit relatively low bond strength, and low loading capacity; although these microstructural/properties aspects tend to improve for nanocomposite coatings. Thermal spray coatings are also anisotropic because of the lamellar structure of the coatings; i.e. properties on the plane parallel to the coating surface are significantly different from that on the plane perpendicular to the coating surface. In depositing nanocomposite coatings, there is a possibility of materials degradation and loss of nanostructure because of the high temperature process. However, such technical detriments can be overcome through control of the spray process parameters.

Specific topics of prior review articles related to thermal spray and nanostructured coatings are summarised in Table 2. Thermal spray processes and nanocomposites have been reviewed independently previously. The intent of the current article, therefore, specifically reviews nanocomposite coatings that are deposited by thermal spray processes in relation to their processing, microstructure, and performance. The feedstock materials for thermal spray can be in the form of powder, wire, or liquid, and the ‘Feedstock Materials’ section of this review paper is dedicated to the production and characteristics of these materials.

‘The thermal spray deposition: process and microstructure’ section presents the formation mechanisms and the nano/microstructure of nanocomposite coatings manufactured by conventional thermal spray processes as well as modern processes such as cold spray, suspension and solution precursor thermal spray (SPTS). Prominent nanocomposite materials deposited by thermal spray are assessed in the ‘Nanocomposite coatings’ section, with emphasis on process dependent phase transformations, and their correlations with nano/microstructure and properties of the coatings. The properties and applications of selected nanocomposite coatings with a view towards future potentials are presented in the ‘Properties and potential applications’ section.

Agglomerated powder

Nanostructured powder can be manufactured through three main routes: (i) mechanical processing, (ii) chemical processing, and (iii) physical or thermal processing. Mechanical processing involves mechanically crushing or milling large particles to nanoscale. Chemical processing involves nucleating or precipitating nanoparticles from solution or slurry, with or without the occurrence of a chemical reaction. On the other hand, thermal processing involves condensation of nanoparticles through rapid cooling of a supersaturated vapour. The examples, advantages and disadvantages of each of these techniques are summarised in Table 3.

As-collected nanoparticles are not ideal as thermal spray feedstock because of the following technical concerns: (i) handling difficulty and health concerns; (ii) not free-flowing and a tendency to aggregate, resulting in pulsed feeding and/or clogging of the injection line and nozzle; and (iii) insufficient inertia and momentum to penetrate the core of the thermal spray jet and establish an optimum trajectory. Injection of nanoparticles into the jet core by boosting the carrier gas flowrate has the unintended consequence of destabilising the thermal spray jet, as the injection force needs to be similar to that imparted by the jet flow. ⁷⁵

The universal solution to this impediment is to agglomerate the nanoparticles using an organic binder. ⁷⁶ Spray drying produces microscopic particles with the ideal spherical morphology for thermal spray feedstock, while still retaining the nanostructure of the original particles. A suspension or slurry, which consists of dispersed nanoparticles in an organic binder, is atomised to form micrometre-sized droplets that are subsequently dried using hot air in a closed chamber. ⁷⁷ Spray drying is an appropriate technique to produce nanocomposite powder because of flexibility in selection of the raw materials.

Spray dried particles are usually porous and weakly bonded, thus they do not exhibit adequate cohesive strength to withstand the turbulence of a thermal spray jet. The as-received particles are typically heat treated or densified using a flame or plasma process. ⁷⁸ It is critical that the post-treatment is optimised to improve the cohesive strength of the particles and, simultaneously, retain their nanostructure characteristics. Figure 7 shows the morphology of a nanostructured yttria-stabilised zirconia (YSZ) particle (Nanox S4007, Inframat Corp., Farmington, CT, USA) agglomerated by spray drying. ⁷⁹ Commercial nanocomposite powder produced via this route is often preceded by a low-cost milling operation to mechanically crush the particles to nano-size, and to uniformly mix different components of the composite together. Specific properties, such as toughness and ductility, can be imparted to the material by encapsulating the spray dried nanocomposite core into a thin (1–3 μm) metal coating to form a duplex structure (PComP, MesoCoat Inc., Euclid, OH, USA). ^80,81

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Figure 7.

a Nanostructured yttria-stabilized zirconia (YSZ) particle agglomerated by spray drying and b higher magnification of the same particle highlighting the individual nano-sized YSZ particles. Reproduced from Lima et al. ³⁹⁴ with permission from ASM International

Another approach that has been used to consolidate nanoparticles for thermal spray feedstock is through suspension plasma spray (SPS). ^78,82 Nanoparticle consolidation using a plasma source involves injecting a suspension of nanoparticles into a radio frequency (RF) plasma using a probe atomiser. The solvent vaporises as the suspension passes through the plasma source and the particles are collected using a series of filters. This approach combines processing steps such as drying, calcination, spray drying, spheroidisation and sieving into a single operation. Figure 8 shows the nanostructure of 40 wt-% zirconia–hydroxyapatite nanocomposite powder before and after consolidation by SPS.

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Figure 8.

Transmission electron micrograph of zirconia–hydroxyapatite (40 wt-%ZrO₂/HA) nanocomposite powder manufactured by suspension plasma spray (SPS). Reproduced from Kumar et al. ⁸² with permission from Elsevier

Nanoparticles or nanotubes can be introduced onto the surfaces of micrometre sized particles, rather than dispersing them within the particles, by mechanical or chemical means to produce nanocomposite feedstock. For instance, multi-walled carbon nanotubes (MWCNTs) were grown on the surface of stainless steel and alumina particles through catalytic CVD. ^83,84 Figure 9 shows the morphology of carbon nanotubes (CNTs) incorporated onto the surface of alumina particles. In the mechanical approach, micrometre-sized particles were blended or milled with CNTs to disperse the nanotubes on the particle surface. ⁸⁵

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Figure 9.

Incorporation of carbon nanotubes onto surface of alumina particles: a before and b after catalytic chemical vapour deposition. Reproduced from Balani et al. ⁸⁴ with permission from Elsevier

Other less favoured methods that have been employed to produce nanocomposite powders for thermal spray feedstock include (i) single step homogeneous mixing of the nano-sized component with matrix material using ultrasonic ⁸⁶ or milling ⁸⁷ techniques, (ii) sol-gel synthesis, ^88,89 and (iii) freeze drying. ⁹⁰ However, these techniques produce particles with angular, irregular, or blocky morphology, a wide particle size range, and attending poor flowability. Figure 10 shows the particle morphology of Al₂O₃–SiO₂ nanocomposite powders produced by sol-gel and freeze drying techniques. Grain growth inhibitors ^91,92 can be introduced into nanocomposite powders to maintain the nanostructure during powder processing that involves heating, such as in sol-gel synthesis and thermal spraying. For instance, boron nitride may be uniformly dispersed as a grain growth inhibitor for WC–Co nanocomposite powders. ⁹³

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Figure 10.

Scanning electron micrographs of Al₂O₃–SiO₂ nanocomposite powder produced by a sol-gel and b freeze-drying techniques. Reproduced from Jiansirisomboon et al. ⁹⁰ with permission from Elsevier

Despite the fact that thermal sprayed nanocomposite coatings have been reported since the late 1990s, ⁴⁶ nanocomposite feedstock is limited commercially. Many reports on nanocomposite coatings have been produced in-house at the laboratory scale. Among the known suppliers of commercial nanocomposite feedstock are Inframat Corp. (Manchester, CT, USA), Millidyne Surface Technology (Tampere, Finland), and MesoCoat Inc. (Euclid, OH, USA). Agglomerated nanocomposite powders that are currently available commercially for thermal spray applications include Al₂O₃–TiO₂, Al₂O₃–TiO₂–CeO₂, Al₂O₃–ZrO₂–SiO₂, Cr₂O₃–SiO₂–TiO₂, Cr₂O₃–TiO₂, SiN–NiCr, TiN–NiCr, TiN–CoCrNiMo, WC–Co, WC–Co–Cr, WC–Ni, WC–Co–NiCr, WC–Co–NiCrFeSiBr, WC–Co–NiCrFe, ZrO₂–MgO, ZrO₂–Y₂O₃, and ZrO₂–Y₂O₃–NiO.

Nano-cored wire

Despite being the preferred form of feedstock material because of their lower costs; ease of material storage and handling; high deposition rate; and high deposition efficiency, nanocomposite wire feedstock has little mention in the literature. Nanocomposite wires cannot be manufactured using the cold drawing or rolling techniques used for conventional wires, because these are limited to malleable and electrically conductive materials. Instead, nanocomposite wires are produced as cored wires, where a metallic sheath surrounds the core of nanoparticle agglomerates. ^94,95 Figure 11 shows the cross-section and nanoparticle distribution in a nanocomposite cored wire. ⁹⁶

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Figure 11.

a Cross-section of nanocomposite cored wire (Praxair Surface Technologies Inc., Indianapolis, IN, USA) and b high magnification SEM micrograph showing the homogeneous distribution of nanocomposite particles within the cored wire. Reproduced from Georgieva et al. ⁹⁶ with permission from ASM International

There are several variations of cored nanocomposite wires, characterised by the type and structure of the core filling and metal sheath materials. The construction of the simplest and most common cored wire is shown in Fig. 12a . Such cored wires are produced by first shaping a metal strip into a U-shaped tube. Pre-agglomerated nanocomposite particles are fed to the interior of the tube. The tube is then closed around the agglomerates to form a sheath and pulled through a die to seal the wire and reduce its diameter to the desired size. ⁹⁷ The nanocomposite particles and agglomerates can be prepared using the techniques discussed in the ‘Agglomerated powder’ section.

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Figure 12.

Schematics of various embodiments of nanocomposite cored wires: a cored wire with a metallic sheath around a core filling of nanoparticles agglomerates, b cored wire with a metallic sheath around a core filling of micrometre-sized metallic particles dispersed in nanoparticles agglomerates, c cored wire with bilayer metallic sheath around a core filling of nanoparticles agglomerates, and d cored wire with a metallic sheath around a core filling of nanoparticles agglomerates and a metallic wire extending through the core filling. Redrawn from Ref. 97

One variation of the nanocomposite wire incorporates micrometre-sized metal particles into the agglomerates of nanoparticles; Fig. 12b . The metal particles will melt with the metal sheath during spraying and act as a binder for the nanocomposite agglomerates. A second, inner layer of metal sheath can also be introduced into the cored wire, Fig. 12c , to serve as a high content alloying element and produce exothermal heat to improve particle flattening. A metal wire that extends through the core filling, Fig. 12d , can serve as an alloying element and improve the agglomerate uniformity during spraying. The incorporation of a metal wire also minimises the amount of oxidised inclusions in coatings.

The suppliers of nano-cored wires for thermal spray applications include Praxair Surface Technologies Inc. (Indianapolis, IN, USA), Inframat Corp. (Manchester, CT, USA), and the NanoSteel Company Inc. (Providence, RI, USA). Inframat Corp. has introduced a range of nanocomposite cored wires known as NanoCore, which includes core fillings such as WC–Co/Inconel 625, WC–Co/NiCr, WC–Co/Stainless steel. The nanocomposite cored wires by Praxair Surface Technologies Inc. (140MXC) and the NanoSteel Company Inc. (SHS 7570, SHS 8000, SHS 9172) are iron-based amorphous alloys that form nanocomposite coatings upon spraying. The compositions of these cored wires are specified in Table 4.

Suspensions and solutions

Liquid feedstocks for thermal spray in the form of suspensions and solution precursors have been pioneered by Gitzhofer et al., ⁹⁸ Tikkanen et al., ^99,100 and Karthikeyen et al., ^101,102 in the mid-1990s. Suspensions (slurries or sols) are fine solid particles that are dispersed within a liquid medium, where the fine particles are produced through mechanical, chemical, and/or thermal routes, as discussed in the ‘Agglomerated powder’ section. ‘Solution precursor’ refers to a mixture of constituent chemicals (e.g. inorganic salts or organometallics) dissolved in a solvent, which will react to form the desired material by high temperature reactions and processes in the thermal spray flame or jet. Tables 5 and 6 summarise the formulations of suspension and solution precursor feedstocks, respectively, for thermal spraying of known nanocomposite coatings.

In suspension and SPTS, the feedstock carrier that is conventionally a gas is replaced by a liquid carrier. As such, there are distinctive factors that need to be considered: (i) particle dispersion and stability, (ii) solvent interaction with the flame or jet, (iii) solid loading, (iv) deposition efficiency, (v) precursor concentration, (vi) heat fluxes, and (vii) heat and momentum transfer. This section reviews the fundamental aspects of suspension and solution precursor feedstock preparation from items (i) to (v). More detailed information on the preparation of liquid feedstock and aspects (vi) and (vii) can be found in the review articles by Pawlowski ^103,104 and Fauchais et al. ^75,105

The dispersion and stability of the particles in a suspension feedstock are crucial to achieve homogeneous coatings. Fine particles tend to agglomerate because of their high surface energy even when they are mechanically agitated. Thus, electrostatic, steric, or electrosteric dispersants are used to disperse and stabilise nanoparticle suspensions by establishing long-range repulsive forces. Electrostatic stabilisation is achieved through acquisition of surface charges, or an electric double layer, through mechanisms such as adsorption, dissociation, ionisation, and/or surface reaction when immersed in an aqueous solution; ^106,107 Fig. 13a . On the other hand, steric stabilisation is induced by adsorption or grafting of polymers, or macromolecules onto surfaces; ^106,107 Fig. 13b . The combination of electrostatic and steric effects is known as electrosteric stabilisation and is typically achieved through adsorption of charged polyelectrolytes; Fig. 13c .

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Figure 13.

Schematics representation of particles interaction with a electrostatic stabilisation, b steric stabilisation, and c electrosteric stabilisation

Most dispersants used to prepare thermal spray suspensions are electrosteric. For instance, cationic polyethyleneimine ¹⁰⁸ and anionic ammonium polyacrylate ¹⁰⁹ polyelectrolytes were used as dispersants for Al₂O₃–YSZ suspensions; Table 5. The stability of particle dispersion is determined by the zeta potential, which is the potential between the charged surface and the electrolyte solution. ¹⁰⁶ A zeta potential of at least ±30 mV is required to achieve a stable suspension. ^110,111 Therefore, stability of a suspension depends on the type of dispersant and pH value of the suspension. In suspension systems where both acid and base components are present, such as in WC–Co where WO₃ is a Lewis acid while CoO is a base, achieving stability becomes an intricate process of achieving the equilibrium point between activation of the dispersant and dissolution of WO₃ or CoO. ^112,113

Additional measures to improve particle dispersion, such as milling and sonication, have also been implemented for nanocomposite suspension feedstocks; Table 5. Plasticisers such as polyvinyl alcohol and polyethylene glycol were added to increase suspension viscosity, which will inherently improve the suspension stability. However, it should be noted that the suspension viscosity should be adjusted by the solid loading and amount of dispersant and plasticisers so that it exhibits a minimum value with sufficient stability and shear-thinning behaviour to facilitate suspension feeding. ^75,104

Selection of the solvent in the suspension and solution precursor feedstock, which is typically water, alcohol, or a combination of both, dictates the interaction of the suspension and precursor molecules with the thermal spray flame or jet. Water has a cooling effect while alcohol heats up the flame or jet, because of their different enthalpies of vaporisation; e.g. enthalpy of vaporisation for water and ethanol are 2·26 and 0·84 kJ g⁻¹, respectively. Alcohol also needs a greater injection velocity than water to penetrate the flame or jet because of its lower density; ¹⁰³ i.e. densities of ethanol and water are 0·79 and 1·0 g cm⁻³, respectively. The heat required to vaporise water is three times that of ethanol ¹¹⁴ while the deposition efficiency of an ethanol-based YSZ suspension is double that of a water-based YSZ suspension. ¹¹⁵

Solid loading is the relative amount of solid particles in a suspension or solution precursor feedstock. Generally, the solid loading of nanocomposite suspensions reported in the literature does not exceed 50 wt-%; Table 5. Typically, the solid loading does not exceed 25 wt-%, except in the case of a sol, because the degree of particle melting decreases with increasing solid loading. ^75,104 Solid loading also influences the particle velocity through its effect on feedstock viscosity; i.e. viscosity increases with increased solid loading. Therefore, solid loading has an important role in the deposition efficiency of the process, where deposition efficiency increases when the particle melting and particle velocity increase. However, it should be kept in mind that particle melting reduces the degree of nanostructure morphology. The deposition efficiency of suspension thermal spray (STS) is about one fourth to one-third of that of conventional coatings; ⁷⁵ i.e. a deposition efficiency of 20% is typical for this type of materials processing, which can be compared to values of 55–80% for conventional thermal spray techniques.

Solution precursor feedstocks are prepared by mixing stoichiometric amounts of inorganic salts or organometallic precursors with water and/or alcohol; Table 6. Important aspects to be considered in the selection of precursor materials are the (i) precursor costs, (ii) product formation and crystallisation temperature, (iii) stability of the precursors, and (iv) potential harmful by-products. Solvents that have low surface tension and low boiling point are recommended to produce dense coatings; e.g. ethanol compared to triethanolamine, as these solvents dissociate and evaporate more easily, which is more likely to lead to full pyrolysis. ¹¹⁶

Unlike suspension feedstock, the solid phase contents of solution feedstocks are often reported in the form of precursor concentration. The resulting solid phase content can be calculated using the volume, concentration, and molecular masses of the respective precursor compounds in the feedstock. A high precursor concentration that is close to the equilibrium saturation concentration is recommended to induce volumetric precipitation for dense coatings. ¹¹⁷ The equilibrium saturation concentration is the concentration at which the rate of dissolution and precipitation of the solid and solution phases, respectively, becomes equivalent. Although the increase of precursor concentration increases the solution viscosity, it has little effect on the surface tension and specific mass, hence the pyrolysis behaviour and crystallisation temperature remain unaffected. ¹¹⁷

Conventional approaches

Formation mechanisms

Thermal sprayed nanostructured coatings via conventional approaches; i.e. deposition of powder feedstock using thermal spray processes other than cold spray, have mostly been deposited using micrometre-sized agglomerates of nanoparticles. Such an approach can be considered as a combination of top-down and bottom-up approaches of producing nanostructured materials because the feedstock are often crushed to nanometre-sized from micrometre-sized particles (top-down), then agglomerated to micrometre-sized and deposited to build-up nanostructured coatings with a thickness in the micrometre range (bottom-up).

Upon entering the thermal spray source, particle melting initiates from the surface of the particle. As agglomerated particles are porous and have low apparent density, the thermal transport properties of these particles are different than those of the conventional, dense feedstock particles. The heat flux from the thermal spray source is retarded by the pores and heat diffusion is lower compared to dense particles of the same diameter. ¹¹⁸ This effect allows partial melting of the particles for effective deposition while retaining the nanostructural characteristics of the feedstock. As such, the particle momentum has to be adjusted by fine-tuning the spray parameters to allow for an appropriate degree of particle melting during in-flight. Large feedstock particles are favoured when sprayed through this route to avoid a high degree of melting. ¹¹⁹ It is stressed that the thermal spray parameters that are employed for conventional feedstocks cannot be used for feedstock of nano-character because any nano-features would be lost during the process.

It should be emphasised that partial melting of feedstock is prevalent in all thermal spray processes and it is not uncommon to ascertain ‘unmelts’ or ‘partial melts’ within a thermal spray microstructure. However, for the thermal spray of nanocomposites, partial and incomplete melting is a preferred mechanism of coating formation because this allows retention of the nanostructure within the feedstock.

Upon impact on the substrate, the molten fraction of the particles solidifies to micrometre-sized artefacts that enhance the coating's bond strength; while the non-molten nanoparticles retain the nanostructure character of the deposit. Therefore, partial melting of the particles causes the loss of nanostructural characteristics to some extent. Nanocomposite coatings sprayed through this approach will not be completely nanostructured, but instead have a bimodal structure, ^75,120 where the nano-sized grains are embedded within micrometre-sized grains. A loss of nanostructure within the molten phase is also contributed by the sintering of the liquid phase, because the sintering time is reduced to the order of nanoseconds compared to solid phase sintering that may take hours. ¹¹⁸ Porosity in the feedstock is introduced into the coatings as porous nanozones when the molten region does not fully infiltrate into the non-molten core of the partially molten particles, which generates improved properties for applications such as thermal barrier coatings and abradable seals. The formation of a bimodal structure through (i) partial melting, and (ii) deposition of agglomerated nanocomposite particles is shown schematically in Fig. 14.

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Figure 14.

Schematic representation of agglomerated nanostructured particle melting and deposition by conventional thermal spray processes (Δ: component 1; ▄: component 2; ○: pore)

Bimodal nano- and microstructure

The microstructure of plasma or HVOF sprayed nanocomposite coatings, as viewed at magnifications of less than 2 000×, resembles the lamellae of conventional thermal spray coatings. However, magnifications greater than 2 000× reveal the nanostructural characteristics of the coatings that are retained through partially or non-molten particles; thus, contributing to the bimodal structured coatings discussed in the previous section; Fig. 15. An example of microstructural comparison between conventional and nanostructured composite coatings is depicted in Fig. 16. The lamellar structure is detected in the molten region of the coatings (marked as ‘A’) while the partially molten regions exhibit embedded particulates (marked as ‘B’ and ‘C’). The extent of the nanostructure in such coatings is determined by the amount of partially or non-molten particles retained in the coatings.

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Figure 15.

Schematic of nanocomposite coatings microstructure deposited by conventional thermal spray processes using agglomerated powder feedstock

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Figure 16.

Cross-sections of plasma sprayed Al₂O₃–13wt-%TiO₂ coatings: a Conventional coatings showing columnar grain boundaries, b nanostructured coatings consisting particulates embedded in partially molten regions, and c high magnification of the region highlighted in b. Reproduced from Goberman et al. ¹²⁵ with permission from Elsevier

Plasma sprayed ZrO₂–Y₂O₃ and Al₂O₃–TiO₂ coatings have exhibited up to 52 and 25% of nanoparticles retainment via the non-molten phase, respectively. ^120,121 Molten splats may consist of strongly bound nano-sized crystallites created during the rapid solidification process. ¹²² Such molten splats have a nano-width columnar structure. It should be noted that a high percentage of nanozones does not necessarily indicate superior performance. There is consensus that a range of nanostructure content is necessary to achieve optimal coating properties; above and below which the specified properties will deteriorate. ⁷⁹ For instance, plasma sprayed Al₂O₃–8wt-%TiO₂ coatings exhibit optimal wear properties at 25% retainment of nanostructure ^123–125 while Al₂O₃–13wt-%TiO₂ coatings perform optimally at 15–20% of nanostructure retainment. ¹²¹

An important microstructural aspect of nanostructured coatings deposited through the conventional approaches is the density of the nanozones. The density of the coatings determines their prospective applications; e.g. porous ZrO₂–7wt-%Y₂O₃ coatings for thermal barrier applications and dense Al₂O₃–13wt-%TiO₂ coatings for wear resistant coatings; Fig. 17. Coatings with specific nanozone density can be engineered through optimisation of particle temperature, velocity, and size distribution. The porous nature of the agglomerated feedstock is preserved in coatings when molten portions of the particle do not penetrate into the porous regions during in-flight, impact, and solidification; whereas a dense structure is achieved by enabling such infiltration.

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Figure 17.

Cross-sections of plasma sprayed coatings with different nanozones density: a porous ZrO₂–7wt-%Y₂O₃ (reproduced from Lima and Marple ¹¹⁹ with permission from Elsevier) and b dense Al₂O₃–13wt-%TiO₂ (reproduced from Fauchais et al. ⁷⁵ with permission from IOP Publishing)

Therefore, porous nanozones are produced by using large agglomerates and achieving a surface temperature that is around the melting point of the material, while dense nanozones are obtained by attaining a surface temperature that is significantly higher than the melting point of the material. ⁷⁵ For instance, 35% porous nanozones were embedded in ZrO₂–7wt-%Y₂O₃ coatings sprayed at an average particle temperature and velocity of 2 670°C and 210 m s⁻¹, respectively. ¹¹⁹ High velocity oxygen fuel is a more effective process compared to plasma spray to produce dense nanostructured coatings because of its high particle velocity. However, the lower particle temperature and shorter in-flight time must be compensated by using smaller particles with a narrow size distribution to provide a high degree of melting so that the agglomerate porosity can be infiltrated.

Cold spray

Formation mechanisms

Despite the fact that nanocomposite coatings deposited by cold spray used similar feedstock as the conventional approaches, the coatings exhibited different microstructures and properties because of the differing deposition mechanisms between cold spray and other thermal spray processes. That is, cold spray operates at a much lower temperature compared to other thermal spray processes. For example, cold spray operates at a maximum of 800°C but typically 400°C, which can be compared to >2 500°C for a plasma spray process. ^52,126 Therefore, deposition via cold spray is based on particle deformation rather than the melting of particles that is experienced for conventional thermal spray processes. The particle velocity must exceed a certain critical velocity that is unique to the material for effective deposition to occur. ^127,128 Particles impacting at velocity below the critical velocity will erode the substrate material without effective deposition. A deformable phase, which is typically a ductile metal phase, is required for deposition of brittle or hard materials such as ceramics so that a composite coating may evolve.

Upon impact with sufficient energy, the cold spray particles undergo high plastic strain rates that result in adiabatic heating of the contact zone surroundings and the material exhibits localised softening. However, the heat from adiabatic shear instabilities is typically insufficient to cause particle melting. Because the particles are deposited without undergoing melting, the nanostructure of the original feedstock is preserved in the coatings. Grain coarsening that typically occurs during droplet formation in the liquid state in conventional approaches is also prevented during the cold spray deposition process. More information on the gas dynamics, deposition mechanisms, and influence of spray parameters by the cold spray method can be found in Refs. 52, 126, and 129. Figure 18 illustrates the formation of nanocomposite coatings via the cold spray route.

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Figure 18.

Schematic representation of nanocomposite deposition via the cold spray process

Coating microstructures

Generally, the nanocomposite coatings deposited by cold spraying agglomerated particles exhibit a dense and crack-free microstructure, with a uniform distribution of particulates among the binder phase; Fig. 19. The lamellar structure that is typical of thermal spray coatings is not distinguishable. ^130–132 The dense coating microstructure compared to the porous agglomerated feedstock is attributed to effective deposition and densification through particle deformation at high impact velocity. The supersonic impact velocity also caused the formation of irregular substrate surfaces with valley and re-entrant features that are effectively filled by the incoming particles. ¹³³ The low porosity, absence of cracks, indistinguishable particle interfaces, and absence of spallation and delamination indicated that the coatings demonstrated good cohesion and adhesion. ^131,134

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Figure 19.

Schematic of nanocomposite coatings microstructure deposited by cold spray processes using agglomerated powder feedstock

As cold spray operates at a much lower temperature compared to the conventional approaches, there was no evidence of particle melting in the coatings and the nanostructure of the feedstock is retained in the coatings. Figure 20 depicts an example of a cold sprayed nanocomposite coating cross-section, and the view of the retained nanostructure at higher magnification. Cold sprayed nanocomposite coatings are more dense and hard compared to their micrometre-structured counterparts because higher particle velocity can be achieved for smaller particles using the same gas flowrate. ¹³⁵ In addition, there is a tendency for the ductile matrix phase to undergo strain hardening during the cold spray process. ¹³⁶

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Figure 20.

Cross-sections of cold sprayed NiCrAl–40vol.-%BN nanocomposite coatings. The arrows indicate cubic boron nitride dispersoids. Reproduced from Luo and Li ¹³² with permission from Springer Science and Business Media

Building up of composites with micrometer-sized hard ceramic particles such as oxides or carbides by cold spray has been difficult compared to plasma and HVOF deposition because of poor deposition efficiency. However, cold sprayed coatings are favoured over conventional methods because the materials do not undergo high temperature phase transitions into undesirable phases. The use of nanostructured feedstock significantly improves the deposition efficiency and density of cold sprayed coatings of such materials; e.g. tungsten carbide–cobalt, WC–Co. ¹³⁷ Such an outcome is explained by the fact that the critical velocity is lowered and there was a higher surface area for contact between the binder phase and hard ceramic particles for effective deposition of nanostructured particles. Cold spraying of conventional feedstock using similar processing parameters result in erosion and/or grain refinement rather than deposition because of insufficient contact between the hard phase and binder.

Despite the fact that the deformation mechanisms of cold spray deposition is advantageous in producing dense coatings, the extent of particle deformation becomes a precarious matter when the preservation of the initial particle structure is critical to maintain effective functional properties of the coatings; e.g. CNTs composites. Similar to other cold sprayed coatings, the nanocomposite coatings incorporating CNTs underwent significant grain refinement associated with the high strain rate plastic deformation of the process. ^138,139 Although the CNTs were retained and distributed uniformly in the coatings, there was evidence of structural damage through shortening of CNTs by ∼30%. ¹³⁸ Fractured CNT tips were also evident as a result of impact and/or shear by incoming particles. However, the relative degree of structural damage, measured by the ratio of D- to G-band intensity (I _D/I _G) in Raman spectroscopy, indicated that the majority of the CNT structural damage occurred during the milling process of feedstock preparation rather than at the deposition stage. ¹³⁹ The extent of structural damage can be reduced by cold spraying under low pressure (0·5–0·6 MPa) conditions. ¹³⁹

Suspension and solution precursor thermal spray

Formation mechanisms

The development of STS and SPTS has rendered thermal spray processes more versatile in the deposition of nanocomposite coatings. (Note: within the literature the acronyms of ‘SPS’ and ‘SPPS’ are prevalent, where the substitution of ‘T’ is replaced by ‘P’ for ‘plasma’. It is considered within the current review that the more generic ‘thermal’ term is appropriate because this encompasses a wider range of process adaptions.) These techniques are able to fabricate nanostructured coatings without the need for costly and time consuming preparation of nanostructured free flowing particles. In addition, STS and SPTS are valuable in depositing composites with immiscible components and thin coatings within the nano- and sub-micron range. The STS method was developed based on the requisite that a denser carrier is needed for the injection and transport of nanoparticles in high enthalpy flow, while the SPTS technique can be considered as a hybrid of chemical and physical methods in producing nanoparticles discussed in the ‘Agglomerated powder’ section. Therefore, the SPTS process can be used to produce nanocomposite coatings with metastable phases. ¹⁴⁰ The thermal spray sources that are typically used for STS are plasma spray and HVOF, although the former is more favoured than the latter, while SPTS typically employs plasma spray.

The STS and SPTS techniques have feeding systems and spray conditions that are distinctive from the conventional powder feeding systems: ⁷⁵ (i) the feedstock carrier is a liquid rather than a gas, (ii) the spray distance is shorter, approximately one-half to one fifth (i.e., 3–7 cm), than in conventional approaches because nanoparticles have low inertia, and (iii) the deposition efficiency is about one fourth to one-third, of that of the conventional approaches, i.e. the DE is typically 15–20% for the STS and SPTS techniques. The feedstock can be injected as a liquid jet or as droplets through atomisation. In mechanical injection, the liquid feedstock is either forced through the nozzle from a pressurised reservoir or through superimposing pressure pulses at variable frequencies using a magnetostrictive rod. ¹⁰⁵ The liquid jet starts to break up on exiting the nozzle because of air resistance. Atomisation uses external forces such as external pressure, acoustic energy, and centrifugal forces, to break up the bulk liquid before directing the atomised particle stream into the thermal spray effluent.

Upon entering the high velocity flame or plasma jet, the suspension droplets may undergo secondary fragmentation or aerodynamic break up into smaller droplets because of shear deformation from the drag force of the droplets. The extent of the aerodynamic break up depends on the Weber number of the droplets relative to the flame or plasma. ¹⁴¹ The solvent encapsulating the nanoparticles is vaporised to form nano- or sub-micron sized aggregates. The aggregates undergo solid phase sintering, where the duration of sintering depends on the flame or jet temperature. Depending on the spray conditions, some of the finer particles are vaporised while others undergo melting and are deposited on impact with the substrate. More details concerning the heat and mass transfer for the STS process can be found in Refs. 75, 104, and 105. The schematics of coating formation mechanisms by STS are depicted in Fig. 21.

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Figure 21.

Schematics of nanocomposite coating formation mechanisms in suspension thermal spray (STS)

In SPTS, the liquid precursor droplets undergo similar aerodynamic break up as in STS on entering the thermal spray stream to form smaller droplets. The aerodynamic break up occurs in the order of microseconds. ¹⁴² The droplets are then subjected to heating and solvent vaporisation, followed by solid phase condensation. The chemical reaction between the precursors is initiated, and the reaction progresses through nucleation and growth until at least one of the precursors is fully consumed. The coating formation mechanisms by SPTS are illustrated in Fig. 22. Depending on the heating rate and nature of the precipitation during heating and vaporisation, different morphologies of dense particles and/or solid shells may be obtained. ¹⁴³ Dense particles are typically formed within small droplets with high solute diffusivity when the solute concentration and precipitation occurs uniformly throughout the droplets, followed by pyrolysis and sintering; Fig. 22a . Some of the particles may interact to form agglomerates and/or large particles. ¹⁰¹

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Figure 22.

Schematics of nanocomposite coating formation mechanisms in solution precursor thermal spray (SPTS)

Supersaturation occurs near the droplet surface when the rate of solvent vaporisation exceeds the rate of solute diffusion such that a high solute concentration builds up near the surface; Fig. 22b . At a critical solute concentration, a porous shell surrounding a liquid core is formed through precipitation. Depending on the permeability of the shells, different routes of internal vaporisation and pressurisation lead to a variation of shell formations. For a low permeability shell, the internal pressurisation internal pressurisation, arising from liquid core vaporisation, leads to shell fracture and the formation of fragmented shells; Fig. 22c . A high permeability shell enables sufficient release of the internal pressure to form an unfragmented, hollow shell; Fig. 22d . Internal pressurisation within an impermeable shell results in shell rupture and secondary atomisation of the liquid core to form disintegrated shells and fine particles; Fig. 22e . Some precursors may form an elastic shell that undergoes inflation from internal vaporisation and pressurisation, which is then followed by deflation and collapse into a precipitate clump; Fig. 22f .

The outcome of these proposed mechanisms is the evolution of particles that exhibit morphologies that then participate in the classical thermal spray deposition process. The dense particles and various shells are melted and deposited on impact with the substrate. The particle morphology depends on the (i) droplet size, (ii) precursors and solute chemistry, (iii) solute diffusivity and solubility, and (iv) droplet thermal history. ¹⁴³ Generally, large droplets and high heating rates result in hollow particles, while small droplets and low heating rates evolve dense particles. ¹⁴³ Further details on the solution — hot gas interaction and numerical aspects of droplet fragmentation, heat exchange, vaporisation rates and internal pressurisation are available in Refs. 75, 104, 105, and 142.

Microstructure: suspension thermal sprayed coatings

Depending on the feedstock characteristics and processing parameters, the microstructure of STS nanocomposite coatings can take the form of a granular nanostructure, ultrafine lamellar structure or a combination of both; Fig. 23. For instance, the microstructure of a Y₂O₃–50vol.-%MgO nanocomposite coating consists of a lamellar structure with entrapped unmelts in three distinct arrangements: (i) a prominent lamellar structure comprising Y₂O₃ (30–140 nm) and MgO (5–30 nm) phases arranged in alternating wavy streaks, (ii) a dispersion of relatively small MgO (30–140 nm) grains in a continuous Y₂O₃ matrix, which resembles precipitation from a supersaturated solid solution, and (iii) unmelted and/or resolidified particles entrapped within lamellae. ¹⁴⁴

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Figure 23.

Schematic of nanocomposite coatings microstructure deposited by suspension thermal spray (STS) process

A granular structure, similar to those of sintered bulk materials, is developed when (i) non-molten or resolidified particles impact with momentum sufficient for deposition, or (ii) non-molten nanoparticulates of one phase are embedded or entrapped within the molten matrix of another phase. For example, the processing parameters to deposit an Al₂O₃–ZrO₂ coating were optimised so that all the Al₂O₃ particles were molten to form a matrix that encapsulated the solid state ZrO₂ nanoparticulates. ¹⁴⁵ On the other hand, a lamellar structure arises from the deposition of molten particles at high impact velocity. However, the lamellar structure deposited by STS may be distinguished from that of conventional thermal spray coatings with respect to their splat size and degree of flattening. For instance, the ultrafine splats observed in suspension plasma sprayed Al₂O₃–ZrO₂ coatings have similar morphology to that of conventional plasma sprayed coatings, but are smaller in diameter; i.e. 1–5 μm compared to 100–150 μm in conventional coatings. ¹⁴⁶

The development of a STS coating microstructure is contingent on the aerodynamic fragmentation of suspension droplets and their penetration into the thermal spray jet or flame. A high degree of droplet fragmentation creates a large dispersion of particle trajectories, which will result in non-uniform particle melting; thereby forming porous coatings with poor cohesion. The extent of droplet fragmentation depends on the (i) characteristics of thermal spray source; e.g. DC vs. RF plasma; (ii) injection parameters; e.g. injection port diameter, injection angle, injection position, and feedrate; (iii) droplet size distribution; and (iv) surface tension of the solvent. ^105,115,147 Further details on the droplet fragmentation and vaporisation mechanisms are available in Refs. 105 and 115. The deformation of suspension particles on impact can take various forms, such as disk-type or splash-type splats, large clusters, overspray, and recondensed particles; which all depend on the droplet degree of fragmentation, thermal history, and kinetic energy. The build-up of these individual deposits collectively determines the final microstructure of STS coatings. The prerequisite and formulations of suspension particles interpreted in relation to the balance of energy are discussed in detail in Ref. 104.

Similar to conventional thermal sprayed coatings, the microstructure of STS coatings is influenced substantially by process parameters that directly affect the particle temperature and velocity. ¹⁴⁸ One such parameter that is unique to the STS process is the solid loading. The increase of solid loading effectively (i) increases the enthalpy required to melt the particle as well as (ii) decreases the particle velocity because of an increase in the suspension-to-droplet mass ratio. Therefore, any improvements in terms of particle penetration resulting from an increased solid loading are offset by the decrease of particle melting. ¹⁴⁷ For example, an increase of solid loading of Al₂O₃–25wt-%ZrO₂ suspension from 5 to 20wt-% produced a thicker, but more porous coating because of increased deposition efficiency and reduced splat flattening, respectively. ¹⁰⁹ Non-optimised spray conditions resulted in large, micrometre-sized pores in STS nanocomposite coatings. ¹⁴⁹ The increase of solid loading was also reported to increase the spacing of horizontal cracks; although there were no significant changes in the vertical crack density. ¹⁴⁸

Another critical microstructural aspect of STS coatings is the inclusion of recondensed particles. Fine particles that were vaporised from the thermal spray jet tend to recondense downstream and deposit on to the substrate or coating surface. The adherence of such particles on the substrate or coating surface, in addition to the non-melted particles that travelled at the periphery of the jet, contribute to defects between two successive passes. Figure 24 shows an example of microstructural defects in a STS coating that has evolved from recondensed particles. Such phenomenon is especially severe for fine ceramic particles that experience high heat flux under STS conditions because they have the tendency to adhere to coating surfaces over 700–800°C. ^105,150 The formation of such defects can be minimised by optimising the suspension mass load and spray conditions.

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Figure 24.

Cross-sections of suspension plasma sprayed ZrO₂–7wt-%Y₂O₃ coating deposited in four passes. Reproduced from Fauchais et al. ¹⁰⁵ with permission from Springer Science and Business Media

Microstructure: solution precursor thermal spray coatings

Contrary to nanocomposite coatings deposited by the previously discussed thermal spray processes, SPTS coatings do not exhibit distinguishable splat boundaries despite the presence of large splat-like deposits; ^151,152 Depending on the spray conditions, SPTS coatings may comprise microstructural artefacts that result from the various stages of chemical reaction and thermal history of precipitates in the thermal spray jet; such as molten splats, hollow splats, spherical particles, fragmented shells, and unpyrolyzed particles; Fig. 25. Therefore, the microstructure would depend on the form of the particle just before impact; e.g. precipitated and molten particles form splats while supersaturated and unpyrolyzed particles form hollow or fragmented shells.

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Figure 25.

Schematic of nanocomposite coatings microstructure deposited by solution precursor thermal spray (SPTS) process

Similar to STS coatings, the microstructure of SPTS coatings also relies substantially on the extent of droplet fragmentation and penetration into the thermal spray jet. The factors influencing the degree of droplet fragmentation in SPTS are similar to those of STS, as described in the previous section. Generally, droplets below 5 μm will result in molten solid particles while larger particles are fragmented and deposited as unpyrolyzed particles. ¹⁰⁵ Xie et al. conducted studies on the deposition mechanisms of SPTS coatings by studying deposits at different regions using a stationary torch ¹⁵³ and single linescan arrangement. ¹⁵⁴ Figure 26 depicts deposits that evolve from different degrees of penetration into the plasma jet.

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Figure 26.

Solution precursor plasma sprayed ZrO₂–7wt-%Y₂O₃ deposits: a Schematic of the temperature variations in plasma jet: (I) the periphery or cold region, (II) the moderately hot region, and (III) the core or hot region; b–e the corresponding deposits formed from the precursor droplets injected into different regions of the plasma jet. Reproduced from Xie et al. ¹⁵⁴ with permission from Springer Science and Business Media

The precursor droplets that travelled in the periphery of the plasma jet (region I in Fig. 26a ) inflated, ruptured, and/or pyrolysed on impact with the substrate to form flakes, irregular agglomerates, or hollow shells; Fig. 26b . Thin films or fine spherical particles, Fig. 26c and d , were formed from the vaporisation and decomposition of precursor droplets that travelled in the moderately hot region (region II in Fig. 26a ). Precursor droplets that travelled in the plasma core (region III in Fig. 26a ) were melted and then solidified to form ultrafine splats; Fig. 26e . Deposition of particles from regions I and II will result in porous coatings, while those from region III result in dense coatings.

The solution concentration has a significant role in determining the rate of volumetric precipitation for dense particle formation in the SPTS process. Compared to spray pyrolysis, the rate of diffusion in SPTS is low compared to the rate of solute concentration increase through rapid solvent evaporation at the droplet's surface. ^143,155 Hence, the volumetric precipitation in SPTS typically initiates at the droplet's surface because the surface has a much higher concentration than the droplet's core. ¹⁴³ Consequently, solutions with high concentration will result in uniform precipitation and the formation of dense particles, hence deposition of high density coatings. On the other hand, solutions with a low concentration will lead to porous coatings because of the formation of shells and hollow particles as volumetric precipitation is less likely to propagate to the droplet's core; Fig. 27.

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Figure 27.

Schematics of droplets evolution in plasma jet at a low solution concentration and b high solution concentration. Solution precursor plasma sprayed ZrO₂–7wt-%Y₂O₃ deposits: c and d are single scan band centre deposits collected at low and high solution concentrations, respectively; e and f are cross-sections of coatings deposited at low and high solution concentrations, respectively. Reproduced from Chen et al. ¹¹⁷ with permission from Elsevier

Low plasma power during SPTS causes incomplete chemical reaction and particle formation, thus the development of porous coatings. ¹⁵⁶ On the contrary, a high plasma power will yield relatively dense coatings because of the increase of thermal energy that promotes pyrolysis and particle melting. However, an increase in plasma power may also cause the formation of cracks and pores because of significant splashing or splat disintegration. Such microstructures also arose from the formation of hollow particles because of high evaporation rates at the droplet's surface and rapid surface solidification at high plasma power. The mechanism of hollow particle formation at high power is similar to that determined for low solution concentrations, where the retained precursor within the hollow particles evaporates or pyrolyses upon impact with the substrate. Therefore, porous regions within coatings can be obtained by optimising the spray conditions to retain more particles with unpyrolyzed precursor, which will then contract to form pores upon pyrolysis. ^151,157

Materials and process comparison

From a general microstructural perspective, coatings deposited by spraying agglomerated nanocomposite particles through conventional approaches (plasma spray, HVOF, and detonation spray) reveal a mixture of fully molten splats, partially molten splats, and non-molten particles; while coatings deposited by cold spray consist of nanoparticles encapsulated within the deformed binder phase. Suspension thermal spray coatings are composed of dense nano-sized spherical particles and very fine splats that result from particle melting. SPTS coatings constitute a combination of flattened splats, hollow splats, spherical particles, unpyrolyzed particles, and fragmented shells, which result from the various routes of in-flight particle formation.

The versatility of thermal spray processes in depositing nanocomposite coatings, regardless of the form of feedstock materials, can be seen in the ‘Conventional approaches’, ‘Cold spray’, and ‘Suspension and solution precursor thermal spray’ sections. A wide range of nanocomposite materials, from ceramic matrix to metallic and polymer matrices, were deposited using conventional plasma spray, HVOF, cold spray, STS, and SPTS processes. The thermal spray processes used to deposit nanocomposite materials are compared in Table 11 based on the analysis of published work. It can be deduced from Table 11 that most of the materials reported are specific to the deposition method, with the exception of several commonly used materials; e.g. Al₂O₃–MWCNT, Al₂O₃–Ni, Al₂O₃–(ZrO₂–Y₂O₃), AlSi–MWCNT, FeCrB–CrB₂, NiO–(ZrO₂–Y₂O₃), TiO₂–HA, WC–Co, and ZrO₂–Y₂O₃.

Each process has specific material requirements because of the feedstock characteristics and variations in process temperature and velocity. In general, conventional flame spray and HVOF processes are employed to deposit cermets of polymer matrix materials such as Fe–epoxy, Nylon-nanodiamond, Nylon–SiO₂, Al₂O₃–NiCr, and TiB₂–NiCrSiB, because of their relatively low processing temperature. Conventional plasma spray, which has a jet temperature of up to 15 000 K, is mostly used to deposit ceramic nanocomposites such as Al₂O₃–TiO₂, Al₂O₃–ZrO₂, Y₂O₃–MgO, and ZrO₂–Y₂O₃. Detonation spray was also used to deposit ceramic nanocomposite coatings; e.g. Al₂O₃–TiO₂ and WC–Co. ^159–161 On the other hand, wire arc sprayed nanocomposite coatings have been limited by the commercial availability of nanostructured feedstock, and are typically formed from metal alloy cored wire that forms carbides, borides, and/or oxides through in-flight reaction and embedment within the metallic matrix. ^96,162–166

Cold spray is typically used to deposit composites with a ductile phase to encapsulate ceramic particulates because of its high particle velocity and the lack of a molten phase; e.g. Al–Al₂O₃, Al–B₄C, HA–PEEK, and TiB₂–Cu. The STS process is mostly used to deposit ceramic nanocomposites that can form stable suspensions with the aid of dispersants and/or plasticisers, for instance, Al₂O₃–ZrO₂, NiO–ZrO₂, TiO₂–HA, and Y₂O₃–MgO. On the other hand, SPTS is used to deposit materials that can be synthesised from precursors to form nanocomposites in-flight; e.g. Al₂O₃–Ni, NiO–ZrO₂, Si–TiO₂, TiO₂–HA, and ZrO₂–Y₂O₃.

Al₂O₃–TiO₂

Thermal sprayed alumina–titania coatings are typically applied to structural materials or machinery for protection against abrasive and erosive wear, cracking, spallation, or corrosion. ^{53,121,124,167} Nanocomposite Al₂O₃–TiO₂ coatings have demonstrated improved wear resistance by up to eight times compared to its conventional counterparts. ^167–170 The improvement in the microstructure and properties of this composite stems from the fact that addition of TiO₂ to Al₂O₃ lowers the overall melting point of the composite; thus lowering the particle viscosity and interlamellar contact of the splats compared to individual ceramic coatings. ^167,168 Various compositions of Al₂O₃–TiO₂ nanocomposite coatings have been reported, including Al₂O₃–3wt-%TiO₂, ¹⁷¹ Al₂O₃–8wt-%TiO₂, ¹²¹ and Al₂O₃–13wt-%TiO₂. ¹⁶⁹ However, the majority of the studies of Al₂O₃–TiO₂ coatings have concentrated on Al₂O₃–13wt-%TiO₂, since this composition was shown to be optimal for wear resistance. ¹⁶⁸

Nano-Al₂O₃–13wt-%TiO₂ coatings have been mainly deposited by plasma spray, since ceramic coatings and their composites are conventionally deposited by plasma spray because of its high temperature. However, HVOF ¹⁶⁷ and detonation spray ^159,160 have also been used to deposit Al₂O₃–13wt-%TiO₂ nanocomposites to create a variation of microstructure and phases that present significant changes in mechanical and tribological performance of the coatings. Conventional commercial coatings deposited using Metco 130 powder, ¹⁷² which exhibit splat quenched single phase γ-Al₂O₃ with clear splat boundaries and contrast between splats as a result of chemical inhomogeneity, ^168,173 are almost exclusively used as the benchmark in studies of nano-Al₂O₃–13wt-%TiO₂ coatings.

Alumina–titania, Al₂O₃–13wt-%TiO₂, nanocomposite coatings were deposited by Gell et al. ^123–125 via plasma spray of spray dried and heat treated nanoparticles. The coatings exhibited two distinct regions: (i) the fully molten regions that consisted of splat quenched γ-Al₂O₃ in nano- and sub-micrometer-sized columnar grains, and (ii) the partially molten regions that consisted of liquid phase sintered α-Al₂O₃ nanoparticles embedded in γ-Al₂O₃ matrix. Subsequent work on plasma sprayed Al₂O₃–13wt-%TiO₂ nanocomposite coatings have similar findings. ^168,173,174 Microstructural comparisons between conventional and nanostructured Al₂O₃–13wt-%TiO₂ coatings are shown in Fig. 16. Metastable γ-Al₂O₃ is formed from fully molten particles because of the high cooling rate, while non-molten or partially molten particles retained the stable α-Al₂O₃ phase from the original feedstock. ^125,169

Although all the reports agreed that TiO₂ phase does not appear on analysis of these coatings, there were different views on the final phase of TiO₂ within the coatings: i.e. (i) Ti⁴⁺ dissolved in γ-Al₂O₃ to form a solid solution; ¹²⁴ (ii) TiO₂ melted preferentially because of its lower melting point, with limited dissolution of α-Al₂O₃ into liquid TiO₂; ¹⁷⁵ and (iii) the formation of non-equilibrium χ-Al₂O₃ TiO₂ phase where the Ti ions randomly occupy the Al³⁺ lattice sites in γ-Al₂O₃. ^123,171 The inconsistencies concerning the TiO₂ phases in coatings could arise because of the variation of volume fraction of TiO₂ in the initial feedstock; which, according to the phase diagram, could alter the solubility of TiO₂ in Al₂O₃.

The variation of microstructure and phases of plasma sprayed Al₂O₃–13wt-%TiO₂ as a function of critical plasma spray parameters (CPSP) has been studied extensively by Gell and co-workers. ^{123–125,169,175} The term CPSP is defined by the ratio of power to primary gas flowrate. Generally, the fractions of γ-Al₂O₃ phase increased with an increase of CPSP until it reached a maximum at CPSP = 390. The γ-Al₂O₃ phase started to decrease with a further increase in CPSP. ^124,125 The increase of CPSP caused an increase of particle temperature, and consequently the increase of volume fractions of the fully molten region and γ-Al₂O₃ phase. Beyond the maximum CPSP, γ-Al₂O₃ phase started to decrease and gave rise to α-Al₂O₃, indicating the occurrence of solid state transformation of γ-Al₂O₃ to α-Al₂O₃ phase beyond a certain particle temperature.

In comparison to the plasma spray system, HVOF deposited particles exhibited a narrow range of in-flight particle temperature and velocity distribution, hence contributing to a more uniform and consistent coating microstructure. ¹⁶⁷ Conventional HVOF coatings were reported to outperform the best plasma sprayed hybrid (i.e. a combination of nano- and sub-micrometre feedstocks) Al₂O₃–13wt-%TiO₂ coatings in terms of abrasion wear resistance and crack propagation resistance. ¹⁶⁷ Nanoparticles achieved approximately 10% higher temperature and velocity than micrometre-ranged particles using the same set of spray parameters because nanoparticles are more porous, have a large surface area, and better capacity to absorb heat. An eight fold improvement in wear resistance was observed in HVOF nanostructured Al₂O₃–13wt-%TiO₂ coatings. ¹⁶⁷ Wear mechanisms of nanostructured coatings will be discussed further in the ‘Tribological properties’ section. The high performance of HVOF coatings compared to plasma sprayed coatings could also be attributed to the variation in phase transformation in both processes. However, no correlation was drawn out in the literature and the phases of HVOF sprayed Al₂O₃–13wt-%TiO₂ coatings were not reported.

Detonation spray has been used to deposit discontinuous Al₂O₃–13wt-%TiO₂ coatings with alternate layers of micro- and nanostructured components, ^159,160 which present an advantage in terms of inhibiting through-thickness crack propagation. The distinctive microstructure of detonation sprayed coatings compared to plasma sprayed coatings is caused by the significant variations of thermal and kinetic states of particles between the two processes. ¹⁵⁹ In detonation spray, there is a natural tendency for segregation of particles with different sizes, where smaller particles are accelerated by the detonation wave and deposited first. Smaller particles also reach a higher temperature and undergo significant melting compared to larger particles.

The thermal and kinetic behaviour of dense and porous particles of differing particle size causes the formation of a bimodal microstructure that is dense near the substrate and loosely bonded towards the surface; Fig. 28. In a fashion similar to plasma sprayed coatings, the nanostructure is retained through deposition of large non-molten or partially molten particles, where TiO₂ melts and forms a liquid sintering phase around unmolten Al₂O₃ particles. ¹⁶⁰ Although γ-Al₂O₃ is still present in detonation sprayed coatings, the amount of γ-Al₂O₃ phase is lesser than α-Al₂O₃ and, in comparison to plasma sprayed coatings, arises as a result of less melting. ¹⁵⁹ The higher particle velocity in detonation spray results in denser coatings with higher hardness compared to plasma sprayed coatings. As well, detonation sprayed Al₂O₃–13wt-%TiO₂ nanocomposite coatings exhibit an increase of 30–60% in hardness compared to conventional coatings. ¹⁵⁹

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Figure 28.

Microstructure of detonation sprayed multi-layer Al₂O₃–13wt-%TiO₂ coating with alternating conventional and nanostructured layers. Reproduced from Cetegen et al. ¹⁶⁰ with permission from Elsevier

Al₂O₃–(Y₂O₃)ZrO₂

The low thermal conductivity of zirconia makes it an attractive material in thermal barrier applications. However, the transition of zirconia between its three polymorphs as a function of temperature causes significant defects when cooled from high temperature. The equilibrium stable phase for zirconia at room temperature is monoclinic; which transitions to tetragonal and cubic phases as the temperature increases. Upon cooling, volume expansion occurs on phase transformations from cubic to tetragonal to monoclinic, which results in stress formation and eventually cracks. Stabilisers, such as yttria, are added to stabilise the tetragonal phase and hinder the transition to monoclinic phase; thus giving rise to YSZ. The two most widely sprayed YSZ compositions are the 7YSZ (ZrO₂–7wt-%Y₂O₃) ^120,122,176 and 8YSZ (ZrO₂–8wt-%Y₂O₃), ^147,177 which are widely used in thermal barrier and solid oxygen fuel cell (SOFC) applications, respectively.

Nano-YSZ coatings have been deposited by conventional, ^120,122,170 suspension, ^147,177,178 and SPPS. ^101,151,179 While plasma sprayed nano-YSZ coatings exhibit a bimodal microstructure, SPPS YSZ coatings were composed of fine splats; namely diameters of 0·5–5 μm with evenly spaced vertical cracks and nano-sized pores, which may be contrasted to a 50–100 μm splat size for conventional plasma spray coatings. The through-thickness vertical microcracks in the SPPS ZrO₂–7wt-%Y₂O₃ coating depicted in Fig. 29 is a unique microstructural artefact introduced to impart strain tolerance in thermal barrier coatings. Such a microstructure contributed to low conductivity, high durability from increased strain tolerance and toughness, and improved thermal cycling life of coatings. ^152,179 The spallation life of the coatings was improved by 2·5 times compared to conventional plasma sprayed coatings of the same bond coat and substrate, although the failure mode of SPPS coatings is similar to that of conventional coatings. ¹⁷⁹

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Figure 29.

Solution precursor plasma sprayed ZrO₂–7wt-%Y₂O₃ coatings: a cross-sectional view, b top view highlighting splat-like feature, and c higher magnification of b highlighting the equiaxed, nano-sized grains. Reproduced from Xie et al. ¹⁵³ with permission from Elsevier

Suspension plasma sprayed YSZ coatings exhibited unmelts, partially molten, and resolidified particles within the coatings, which could be minimised by optimising the heat transfer, particle fragmentation, and particle trajectory through the variation of spray distance, injection velocity, injection location, solid loading, and plasma enthalpy. ¹⁴⁷ At high magnification, the splat-like deposit revealed angular nanometre-sized grains ranging from tens to hundreds of nanometres. ¹⁵³ Despite having more than one constituent, and therefore qualify as a composite, nano-YSZ will not be discussed further because the free yttria content in YSZ is minimal and localised compared to non-transformable tetragonal zirconia that it has not been detected in most studies, ^156,180,181 but was only indexed sporadically by selected area diffraction (SAD). ¹⁵¹ This section will, therefore, focus on YSZ composites with a significant alumina component; i.e. Al₂O₃–ZrO₂.

Alumina was incorporated in YSZ to enhance the abrasion, wear, oxygen diffusion, and thermal resistances of conventional YSZ coatings. ¹⁵⁸ Addition of Al₂O₃ has been evidenced to limit the grain growth of ZrO₂ during heat treatment, where grain growth of 7YSZ is three times greater than Al₂O₃–40wt-%ZrO₂. ¹⁸² The restricted growth of ZrO₂ was attributed to (i) the limited solid solubility of Al₂O₃ and ZrO₂, where the two phases are mutually soluble only after melting; and (ii) the increased diffusion path relative to single phase oxides.

Similar to nano-YSZ coatings, Al₂O₃–ZrO₂ coatings have also been deposited by conventional, ^183,184 suspension ^109,148 and SPPS. ^140,182 The rapid solidification of thermal spray processes is conducive to the extension of the Al₂O₃–ZrO₂ solid solution to form metastable amorphous phases, which could potentially (i) exhibit unique microstructures that would reflect on the electrical, optical, magnetic, thermal and mechanical properties; ^185–187 and (ii) provide a reservoir of chemical species for surface chemical reactions during heat treatment. ¹⁴⁰ The amount of Al₂O₃ incorporated in Al₂O₃–ZrO₂ nanocomposites ranged from 5 to 80 wt-%. ^183,184 Most studies on Al₂O₃–ZrO₂ coatings have examined the near eutectic composition of Al₂O₃–40wt-%ZrO₂, where the eutectic point of the Al₂O₃–ZrO₂ pseudobinary system is 58wt-%Al₂O₃. ^188–190 The near eutectic composition is favoured because the lowered melting point compared to pure Al₂O₃ and ZrO₂ promotes the formation of a nanostructured, metastable solid solution.

Plasma sprayed Al₂O₃–ZrO₂ nanocomposite coatings were characterised by a typical bimodal microstructure composed of nanozones of unmolten and partially molten particles, and a dense molten region. ¹⁹¹ The main difference identified between nanostructured and conventional coatings was the distribution of pores, where nanostructured Al₂O₃–ZrO₂ coatings exhibited a homogeneous pore distribution of equal size compared to the wide and inhomogeneous pore distribution of conventional coatings. ¹⁸³ The difference in homogeneity of pore distribution, in addition to the influence of grain boundaries, contributed towards greater load carrying capacity and less fluctuations of hardness and elastic modulus values in nanocomposite coatings compared to their conventional counterparts. ¹⁸³ The molten regions of coatings revealed the lamellar structure of typical thermal sprayed coatings, ¹⁹² although nanostructured coatings indicated better chimerism between lamellar interfaces compared to conventional coatings because of more homogeneous melting. ¹⁸³

Most studies have reported t-ZrO₂ to be the main phase of as-sprayed coatings with metastable γ-Al₂O₃, although small amounts of m-ZrO₂ ¹⁹² and c-ZrO₂ ¹⁸³ have also been found. Such discrepancies are most likely caused by the differences in composition, initial phase of the feedstock, and/or feedstock processing. The main phase of ZrO₂ in the feedstock is mostly retained in plasma sprayed coatings in the presence of Y₂O₃ and/or Al₂O₃. Yttrium ions are more stable substitutional solutes in ZrO₂ compared to aluminium ions because of their larger ionic radius, thereby making it a more effective stabiliser. However, the use of Al₂O₃ as an additional stabiliser significantly reduced the content of metastable γ-Al₂O₃, which tends to nucleate in favour to α-Al₂O₃ upon solidification by virtue of its low nucleation and interfacial energy for transformation. ^184,193 Addition of other high melting point oxides and carbides; e.g. MgO and SiC, has also been known to give rise to α-Al₂O₃ formation by increasing the amount of available unmolten, nucleation sites. ¹⁸⁴ Such escalation in α-Al₂O₃ content, in addition to the dense nanostructure, contributed to an increase of coating hardness, wear and corrosion resistance. ^184,194

The SPS process has been used to engineer Al₂O₃–ZrO₂ nanocomposite coatings with different architectures; e.g. dense or porous, graded porosity or composition, and finely structured multi-layers. ¹⁹⁵ The varying architecture is achieved by adjusting the interrelated operating parameters: (i) heat and momentum transfer to the particles (related to the plasma gas mixtures), (ii) particle momentum and thermal inertia (related to particle size and spray distance), and (iii) heat flux imparted by the plasma flow to the substrate and previously deposited layers. The plasma auxiliary gas, plasma current, gas flowrate, gas composition, feedrate, particle size distribution, and solid loading contribute to the particle temperature and velocity profiles, which consequently influence phase formation in the coatings. ¹⁴⁸ Low particle velocity (<650 m s⁻¹) favours the formation of stable α-Al₂O₃ and tetragonal ZrO₂ phases while high particle velocity (>750 m s⁻¹) favoured the formation of metastable γ-Al₂O₃ and cubic ZrO₂ phases. ¹⁰⁸

The SPS Al₂O₃–ZrO₂ coatings were dense, with dense spherical particles of ∼100 nm and ultrafine splats that exhibit similar morphology to conventional plasma sprayed coatings but is more refined; i.e. 1–5 μm compared to 100–150 μm in conventional coatings. ¹⁴⁶ Despite having a molten phase, indicated by the presence of ultrafine splats, SPS coatings did not reveal the presence of γ-Al₂O₃ phase that is observed in conventional plasma sprayed nanocomposite coatings, as the metastable phase was transformed to α-Al₂O₃ by the convective flux of the plasma jet at the coating surface because of the short spray distance. ¹⁹⁴ The α-Al₂O₃ and tetragonal ZrO₂ grains were distributed homogeneously in SPS coatings, while the conventional coatings showed distinct boundaries between the immiscible Al₂O₃ and ZrO₂. ¹⁹⁴

Suspension plasma sprayed Al₂O₃–ZrO₂ coatings were highly crystalline and composed of fine lamellae of alternating Al₂O₃ and ZrO₂ layers with segregated crystalline phases. ¹⁵⁸ The segregation arose because of the high affinity of ZrO₂ to the solvent over Al₂O₃, which allows the solvent to entrain ZrO₂ particles as it flows from the core to the droplet surface during evaporation. On the other hand, suspension HVOF sprayed coatings were amorphous and revealed uniformly dispersed ZrO₂ particles in an Al₂O₃–ZrO₂ matrix. ¹⁵⁸ Figures 30a and 30b show the microstructure of granular and lamellar structured Al₂O₃–ZrO₂ coatings deposited by high velocity suspension flame spray and SPS, respectively. The finely dispersed ZrO₂ particles in HVOF sprayed coatings in addition to its larger grain sizes, when compared to SPS coatings, indicate that the particle temperature was lower and the melting was incomplete. As a consequence, SPS coatings exhibit high hardness and high abrasion resistance, while suspension HVOF sprayed coatings show low thermal diffusivity and high erosion resistance.

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Figure 30.

Cross-sections of Al₂O₃–ZrO₂ nanocomposite coatings deposited by a high velocity suspension flame spray, HVSFS (reproduced from Gadow et al. ¹⁴⁵ with permission from Elsevier) and b suspension plasma spray, SPS (reproduced from Tingaud et al. ¹⁰⁹ with permission from Elsevier)

Porous Al₂O₃–40wt-%ZrO₂ nanocomposite coatings were produced by using helium in place of hydrogen as the secondary plasma gas; ¹⁴⁸ Fig. 31. The porosity content was increased from 0% in the crack free regions of the hydrogen-deposited coatings to a maximum of 8% in helium-deposited coatings. The increase of plasma stability with the use of helium, coupled with its high conductivity generates a wider plasma jet core that promotes the deposition efficiency of fine particles. Dense coatings deposited by STS tend to develop microcracks from thermal stress relief upon cooling. This arises because of the short spray distance and greater number of passes needed to attain the required thickness in STS compared to conventional thermal spray. These different thermal spray parameters may generate higher temperature on the coating surface, namely 860°C compared to ∼200°C for conventional thermal spray processes; ¹⁰⁴ Fig. 31a . Coatings with uniformly distributed porosity are desirable in applications such as thermal barrier coatings, abrasive coatings, and the anode layer of solid oxide fuel cells.

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Figure 31.

Suspension plasma sprayed Al₂O₃–40wt-%ZrO₂ coatings deposited using a hydrogen and b helium as secondary gas. c and d are the same coatings as a and b, respectively at 2 000× magnification. ¹⁴⁸

On the contrary, nanocomposite ZrO₂–10mol.-%Al₂O₃ coatings deposited by SPPS exhibited three distinct microstructural features; Fig. 32. The coatings were dominated by (i) a nanostructure (region 1 in Fig. 32), and (ii) a sub-micrometre structure (region 2 in Fig. 32); with (iii) some large spherical grains (region 3 in Fig. 32). ¹⁴⁰ The nanostructured regions were composed of distinct fine (∼10 nm) and coarse (10–40 nm) nanostructured zones that consist mainly of t-ZrO₂ phase, with some m-ZrO₂ grains in the coarse nanostructured zones. The metastable t-ZrO₂ phase was formed as a result of stabilisation by Al³⁺ in the ZrO₂ solid solution. The sub-micrometre regions were composed of 50–100 nm twinned or equiaxed grains of m-ZrO₂ phase. The occasional large spherical grains were primarily orthorhombic zirconia, ‘o-ZrO₂ phase’, ranging from 0·3 to 1 μm. Formation of o-ZrO₂ phase was speculated to arise from aqueous reaction in the mesoporous zirconia-containing precursor. Amorphous aluminium-rich phase was found at the grain boundaries in the form of intergranular pockets.

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Figure 32.

Transmission electron microscope bright-field images of solution precursor plasma sprayed ZrO₂–10mol.-%Al₂O₃ coatings exhibiting three types of structures: (1) nanostructure, (2) sub-micrometre structure, and (3) large spherical grains. Reproduced from Vasiliev et al. ¹⁴⁰ with permission from Elsevier

The presence of amorphous phase surrounding the nanocrystalline grains in plasma sprayed, SPS and SPPS nano-Al₂O₃–(ZrO₂–Y₂O₃) coatings contributed largely towards the decrease in thermal diffusivity, ¹⁹² because of improved phonon scattering and internal boundary scattering. ¹⁹⁶ In addition, nanograins provide a larger grain boundary volume for phonon scattering by virtue of its comparable crystallite size to the mean free path of phonons. Thermophysical properties of thermal sprayed nanocomposite coatings will be discussed further in the ‘Thermophysical properties’ section.

WC–Co

Tungsten carbide, WC, is well known for its exceptional hardness and wear resistance. These hard but brittle carbide particles are frequently bound by a ductile metal matrix, most commonly cobalt, to improve the fracture toughness. Tungsten carbide–cobalt, WC–Co, is hence used extensively in applications that require resistance against abrasion, sliding, fretting and/or erosion; ^52,53 e.g., drills, gears, pistons and valves. Physical properties such as hardness, strength and wear resistance of WC–Co depend largely on the WC grain size and volume fraction. ¹⁹⁷ For thermal spray coatings, such properties are also determined by the porosity and phase composition of the carbide and binder phase. ¹⁹⁷

While the WC grains in cemented carbides tend to form a continuous network of contacting particles within the cobalt matrix, ¹⁹⁸ the carbide particles in thermal spray coatings are more discrete because of the relatively high cobalt content used in thermal spray feedstock and the loss of WC during deposition. ¹⁹⁷ The nominal content of cobalt in thermal spray coatings is typically 12 and 17 wt-% compared to the range of 6–10 wt-% used for sintered materials. ¹⁹⁷ The most commonly deposited WC–Co composite coating is WC–12Co, which has been deposited via HVOF, ¹⁹⁹ detonation spray, ¹⁶¹ cold spray, ^133,200 and SPS processes. ¹¹³ Nanocomposite coatings with nominal compositions of WC–15Co ^201,202 and WC–17Co ^135,137 have also been documented.

Generally, WC–Co tends to undergo a combination of decarburisation, oxidation, reduction, and dissolution reactions when deposited by conventional thermal spray processes, such as plasma spray, HVOF, and detonation spray, because of the high processing temperature. Hard and brittle phases such as W₂C, W, and Co_xW_yC_z, and WO₃ were found in thermal spray WC–Co coatings. ¹⁹⁷ During spraying, cobalt is melted first, followed by the dissolution of carbon from the edges of WC grains into cobalt which leads to decarburisation of WC particles. Upon impact and cooling, W₂C, W, and the remaining WC precipitate from the melt while the remaining liquid forms an amorphous matrix that contains W, C, and Co; ²⁰¹ Fig. 33.

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Figure 33.

Transmission electron micrograph of nanostructured tungsten carbide–cobalt, WC–Co coating deposited by detonation spray, depicting different phases in regions at different melting states. Reproduced from Lim et al. ¹⁶¹ with permission from Elsevier

Nanocomposite WC–Co coatings underwent greater degrees of decomposition and dissolution to form W₂C and W compared to conventional coatings owing to the greater surface area to volume ratio of nano-WC particles. ^161,201 The carbon content decreased by 41% for nanocomposite WC–Co coatings compared to only a 24% reduction in conventional coatings. ²⁰¹ The crystallinity index of nanocomposite WC–Co is also lower than that of conventional coatings. ²⁰¹ Despite having significantly higher deposition efficiency (>70%) ¹⁹⁹ and improved hardness compared to conventional coatings, the greater decomposition to form brittle phases in nanocomposite WC–Co coatings led to their inferior wear resistance. ^198,202 In addition, the fine WC grains in nanostructured coatings resulted in rapid pullout of the hard phase and high wear rates compared to larger WC grains that were more resistant to removal by relatively smaller abrasive particles. ²⁰³

Achieving the right balance of melting would produce good bonding while limiting the degradation and other reactions between constituents in the coating, to produce a more wear resistant matrix phase. This is usually achieved by regulating the process parameters such as the agglomerate size, fuel chemistry, and oxygen–fuel ratio. ²⁰⁴ Another approach is to apply a protective layer over the feedstock particles; e.g. cobalt-coated WC–Co particles, where the thin protective layer effectively shields WC particles from direct exposure to the oxidising flame; thereby minimising the degree of WC decomposition and limiting excessive particle melting. ²⁰⁵

Cold sprayed WC–Co coatings are favoured over the conventionally deposited coatings because the WC particles do not undergo high temperature phase transformation into hard and brittle phases. In fact, nanostructured WC–Co coatings are the most widely reported nanocomposite coatings deposited by cold spray. The coatings did not exhibit any evidence of phase transformations or degradation because the deposition occurred in the solid-state. ^133,135 However, building up of WC–Co coatings by cold spray has been difficult compared to deposition by HVOF and detonation spray because of the poor deposition efficiency of the WC particles.

The use of nanostructured feedstock significantly improves the deposition efficiency and density of cold sprayed WC–Co coatings. ¹³⁷ Crack-free WC–Co coatings with a porosity of <3·8% and thickness up to 900 μm have been deposited by cold spraying agglomerated nanostructured feedstock. ^135,137 The coatings also exhibited a more smooth coating-substrate interface compared to conventional coatings. ¹³⁷ This outcome arises because (i) the critical velocity is lowered and (ii) the contact area between the Co binder and WC particles was greater. Cold spraying of conventional feedstock using similar processing parameters resulted in erosion and/or grain refinement rather than deposition because of insufficient contact between the hard phase and binder. ¹³⁷

On the contrary, SPS produced WC–Co coatings with variable porosity levels from less than 0·2 to 9%. ¹¹³ Coatings with low porosity revealed substantial degradation of WC–Co because of extensive melting of the particles, which caused a significant decrease in hardness. Similar to HVOF coatings of agglomerated WC–Co, the SPS coatings were composed of W₂C, W₃C, and W from the decomposition and reduction of the WC phase, along with the presence of a brittle amorphous cobalt phase.

Distinct carbide degradation during the SPS process is indicated by the decrease in carbon content and low crystallinity. The high particle temperature, in addition to the large surface area of fine carbide particles in the SPS process, increased the surface reactivity of the particles for oxidation and rapid dissolution of WC grains. Thus, the degradation of carbides and binder can be restricted by lowering the particle temperature; e.g. varying the suspension feedrate or increasing the total thermal load.

CNT-reinforced nanocomposites

Carbon nanotubes have evolved as a material with exceptional mechanical, electrical, and thermal properties. ^206–208 Physical properties of CNTs depend on their length, diameter, chirality, and orientations. ^209–212 The high strength of CNTs, namely elastic modulus of ∼0·9 TPa and tensile strength of ∼150 GPa, ^213,214 places them as a reinforcement candidate for composite materials. However, research on CNTs reinforced ceramic and metal matrix composites has been scarce compared to their polymeric counterparts because of the processing constraints in attaining composites with (i) homogeneous CNT dispersion, (ii) effective CNT retention in relation to chemical and structural stability, and (iii) good CNT-matrix interfacial bonding. ²¹⁵

Thermal spray is emerging as a promising processing route to incorporate CNTs into ceramic and metal matrix composites, not only as coatings but also as bulk free standing components. Most CNT-reinforced composites are manufactured using MWCNTs, which consist of many single-walled CNTs that are arranged in a concentric manner. Conventional thermal spray processes, including flame spray, ⁸³ plasma spray, ^85,213 and HVOF, ^215,216 as well as cold spray ^138,139 have been used to deposit CNT-reinforced nanocomposites with ceramic and metal matrices; e.g. Al₂O₃, Al–Si, copper, stainless steel, and hydroxyapatite. Carbon nanotube-reinforced hydroxyapatite coatings will be discussed in the next section, because they require special considerations in relation to toxicology in biomaterial applications.

Free-standing CNT-reinforced Al₂O₃, ^84,85,216 and hypereutectic Al–Si nanocomposites ^217–219 with thicknesses over 1 000 μm were manufactured by HVOF and plasma spray forming. As lightweight CNTs are difficult to retain in the core of a thermal spray jet and are often lost as overspray, improving the feedstock flowability and deposition efficiency by using spray dried powder rather than blended ones enabled the thickness limit to be extended to 5 000 μm, ²²⁰ The different methods of CNT incorporation; i.e. agglomerated vs. surface dispersion, resulted in the variation of in-flight particle characteristics, and consequently the microstructure and properties. ²²¹ Thermal spray formed CNT-reinforced nanocomposites are composed of a two-phase structure: (i) matrix containing uniform CNT distribution, and (ii) CNT-rich clusters. ^215,220,222

The uniformity of CNT distribution in the structure relies on the composition of the feedstock. The degree of CNT clustering increased when the concentration of CNT increased, because of the formation of a CNT mesh through capillary action generated by the surface tension of the molten alloy matrix. ²²⁰ Formation of aluminium carbide, Al₄C₃, which will critically deteriorate the mechanical properties of the nanocomposite, was suppressed during thermal spray because of the thermodynamic stability of the cylindrical graphite-structured CNTs in the aluminium matrix at high temperature. ²¹⁷ The bimodal grain matrix and crack bridging effects of the CNTs led to the improvement of fracture toughness of Al₂O₃–CNT by 43%. ⁸⁵ The mechanical properties of AlSi–CNT were also enhanced by the CNTs, as well as primary silicon particle reinforcement. Plasma spray formed AlSi–MWCNT nanocomposite exhibited a porosity level that was twice that produced by HVOF; thus giving them lower hardness, H = 1·65±0·11 GPa and elastic modulus values, E = 55·5±9·6 GPa compared to HVOF coatings, H = 1·93±0·14 GPa and E = 82·8±9·3 GPa. ²¹⁹ A greater extent of grain refinement of primary silicon particles deposited by HVOF also contributed to the better mechanical properties because they act as potential barriers to dislocation movement. ²²³

Apart from achieving a homogeneous CNT distribution, the CNT-matrix interfaces require low thermal resistance to take advantage of the high thermal conductivity of CNTs, ¹³⁹ as well as good interfacial bonding for the shear stress transfer in short fibre reinforced composites. ¹³⁸ Improvement in the wettability and interfacial adhesion between AlSi and CNTs was achieved through the formation of an ultrathin β-SiC layer at the AlSi–CNT interface, which promoted load transfer to the reinforcement phases. ²¹⁸ Furthermore, CNT inclusions impeded grain growth by restricting grain boundary mobility through grain pinning and the formation of an interfacial amorphous layer. ²²²

Cold sprayed CNT-reinforced composite coatings exhibited homogeneously dispersed CNTs throughout copper ¹³⁹ and Al–Si matrices. ¹³⁸ The solid-state deposition of Cu–CNT coatings resulted in a dense microstructure and ‘clean’ interfaces with no intermediate amorphous compounds; thus exhibiting high thermal diffusivity compared to pure copper coatings. The fibre structure of the CNTs was retained after spraying, but the CNTs were moderately damaged from the impact and shear processes during milling ¹³⁹ and deposition. ¹³⁸ Defects such as kink formation, neck fracture and graphite layer peeling were reported.

The coating surface revealed protruded CNTs that were incorporated in the copper matrix to form a closely attached interface, while the length of CNTs at the cross-sections was maintained at several micrometres and located between elongated copper splats. In the case of AlSi–CNT coatings, the spray dried Al–Si particles that underwent disintegration during cold spray deposition fragmented the CNTs and both components were embedded and mechanically bonded to the aluminium matrix. The coatings exhibited a wide range of elastic modulus values because of inhomogeneous distribution of Si-rich regions as well as porosity that was transferred from the spray dried particles to throughout the coatings.

The manufacture of CNT-reinforced nanocomposites by suspension or SPTS has not yet been reported in the open literature. Nevertheless, the prospects and feasibility of these manufacturing routes are promising considering their capacity to achieve intimate CNT-matrix interfacial interactions at the molecular level of mixing. It is also worth noting that CNT-reinforced nanocomposites have been successfully fabricated through the spray pyrolysis process. ²²⁴

Hydroxyapatite-based nanocomposites

Research on plasma sprayed hydroxyapatite, HA [Ca₁₀(PO₄)₆(OH)₂], has become prominent in biomedical applications since its inception in the 1980s. ^225–227 The structural similarity of HA to bone mineral, which constitutes 60–70 wt-% of natural bone, help promote the fixation and stabilisation at the bone-implant interface. ^228–230 Despite the biocompatibility edge offered by HA, it lacks the mechanical strength to provide the load-bearing capacity required by biomedical implants. Thus, HA is customarily applied as a coating onto metal substrates, such as titanium alloys, to optimise the biological and mechanical attributes of the implants. Bioinert reinforcing materials, such as alumina and titania, are also incorporated to impart strength and toughness to brittle HA without compromising its biocompatibility and bioactivity. Hydroxyapatite has been doped with 5 wt-% silver and then incorporated with a polyether ether ketone (PEEK) matrix to produce nanocomposite coatings that impart substantial anti-bacterial properties to the coatings. ²³¹

Nanostructured hydroxyapatite composite coatings have been demonstrated to not only exhibit superior mechanical performance, but also enhanced biocompatibility with osteoblast cells. ⁸⁷ The improved biocompatibility was attributed to the effect of nanotexture or nanoroughness on the adsorption of proteins such as vitronectin and fibronectin. ^232,233 These proteins mediate adhesion of anchorage-dependent cells such as osteoblasts and, therefore, increase cell reproduction on the surface. Various compositions of HA nanocomposites have been deposited by plasma spray, ^234,235 HVOF, ^236,237 cold spray, ²³¹ SPS, ^238,239 and SPPS; ²⁴⁰ with HA fractions ranging from 10 to over 95 wt-%. The biological and mechanical performance of HA nanocomposite coatings are phase-dependent and, thus, vary with the deposition process.

At high temperature (namely >1 000°C), hydroxyapatite undergoes decomposition to tricalcium phosphate, TCP (Ca₃(PO₄)₂), tetracalcium phosphate, TTCP (Ca₄P₂O₉), and calcium oxide (CaO), in addition to the formation of amorphous phase. ^241,242 Dehydroxylation of HA also occurs to produce oxyhydroxyapatite, OHA and oxyapatite, OAp. ²⁴³ Such phases decrease the purity and crystallinity of HA in the coatings, and consequently result in expedited coating dissolution because these phases have much higher dissolution rates than HA; i.e. ACP≫TTCP>α-TCP>OHA>β-TCP≫HA. ^243,244 Thus, HVOF is the preferred method of depositing HA nanocomposite coatings compared to other conventional thermal spray processes.

HVOF-deposited HA–ZrO₂ nanocomposite coatings exhibited a layered and dense structure, with evenly distributed ZrO₂ particles. ²³⁷ In addition to the high velocity deposition, the distribution of nano-ZrO₂ particles at the HA splat interfaces also contributed to the low porosity. Limited particle melting was related to high crystalline HA phase retention, where the as-sprayed coatings revealed 74% crystallinity. Secondary phases such as α-TCP, t-ZrO₂, and CaZrO₃ were present in the coatings, although these same phases were also found in the feedstock; thereby indicating minimum phase transformation during the HVOF process. The absence of β-TCP, TTCP, and CaO further confirms the limited phase transformation of HA. ²³⁷ The significance of the melting fraction on the phase composition and properties of HA coatings has been discussed by Khor et al. ⁷⁴

The lack of tetragonal to monoclinic ZrO₂ transformation was favourable to improving mechanical properties, and was attributed to the limited heating of nano-ZrO₂, rapid cooling and the formation of residual stress at the HA/ZrO₂ zone during the quenching of molten particles. The presence of CaZrO₃ phase suggests that there has been reaction between HA and ZrO₂ during processing. However, the fraction of CaZrO₃ phase in the coatings only increased 0·2 wt-% from the 1·0 wt-% in the feedstock; indicating that the majority of this phase arose during feedstock consolidation through SPS, rather than the HVOF deposition process.

On the contrary, HA–TiO₂ nanocomposite coatings deposited by HVOF demonstrated significant phase transformation from anatase to rutile TiO₂, where the amount of anatase decreased from 86% in the feedstock to 17% in the coatings. ²³⁰ This phase transition occurred because rutile is the most stable among all the polymorphs of TiO₂. The retention of small amounts of anatase TiO₂ in the partially molten zones greatly enhanced the mechanical and photocatalytic properties of the coatings. ²³⁶ Despite the fact that CaTiO₃ has been found in HVOF HA–TiO₂ coatings with similar compositions, ²⁴⁵ no such phase was detected in the nanocomposite coatings. ^87,230

The absence of CaTiO₃ phase in addition to the lack of thermal decomposition products of HA in the coatings arises because of (i) the insensitivity of XRD or masking by HA peaks owing to the low fraction of HA, i.e. ∼10–20 wt-% compared to 90 wt-% in HA–ZrO₂ coatings, (ii) the short residence time in the HVOF jet, and/or (iii) the feedstock manufacturing process, where they were mechanically blended rather than chemically or thermally consolidated. However, there was some diffusion between HA and TiO₂, which was indicated by the coexistence of titanium, calcium, and phosphorus atoms at microregions between adjacent HA and TiO₂ splats. ⁸⁷ Formation of such products from the chemical reaction between HA and TiO₂ have been claimed to promote apatite growth. ^87,245

In applications requiring photocatalytic performance, nanocrystalline anatase is more favourable as photocatalytic TiO₂ is optimised in this form, ²⁴⁶ where it is most stable at particle sizes <11 nm. ²⁴⁷ While HA requires a low quenching rate to minimise the formation of secondary phases, formation of anatase TiO₂ is favoured by a high quenching rate. A dual injection system that combines conventional powder feeding of HA and SPPS of TiO₂ was used to overcome this exigency to deposit HA–TiO₂ nanocomposite coatings that contained substantial anatase. ²⁴⁰ Despite the benefit of the novel injection system, the coatings maintained the presence of rutile TiO₂ and traces of TCP, TTCP, and amorphous phases. Graded multi-layered HA–TiO₂ coatings deposited by SPS also exhibited similar phases. ²³⁸ Such phases suggest that the anatase to rutile phase transformation at 400–1200°C ²⁴⁸ still occurs regardless of the shorter in-flight time of the particles in SPPS and SPS compared to conventional thermal spray processes.

The deposition of HA nanocomposite coatings by plasma spray, with the incorporation of CNTs, improved coating crystallinity from 53·7 to 80·4%. ²³⁴ The improvement in crystallinity stems from the fact that CNTs act as nucleation sites for the crystallisation of HA and aid the preservation of its inherent crystal structure. ^249,250 During rapid cooling of the plasma spray process, thermal energy is isolated near CNT surfaces surrounded by HA melts because CNTs have three times higher thermal conductivity (∼3×10³ W m⁻¹ K) than HA (∼0·7×10³ W m⁻¹ K). ²³⁴ This allows HA to nucleate over a longer duration in CNT-distributed regions compared to cooling in the absence of CNTs. Despite the improvement, the crystallinity of plasma sprayed HA–CNT coatings is lower compared to that of the feedstock. Transformation of HA to secondary phases was evident in the form of phosphate-rich needle-shaped regions. ²³⁴

While conventional plasma sprayed HA coatings exhibit the typical structure of molten splats, intersplat cracks, micro porosity, and partially fused particles; HA–CNT coatings constitute resolidified fine and nodular HA particles. The nodular particles in the nanocomposite coatings contributed to increased surface roughness, and consequently enhanced cell adhesion and growth. ²³⁵ The distribution and retention of CNTs in HA coatings were similar to that of other CNT reinforced coatings discussed in the ‘CNT-reinforced nanocomposites’ section, where the uniform distribution of CNTs in the coatings indicated maximum retention from the high temperature and velocity deposition. ²³⁴

Wear and fracture toughness were improved by up to 56% with CNT-reinforcement through the anchoring effects of CNTs. ²³⁵ The reinforcing mechanisms of CNTs will be discussed further in the ‘Mechanical properties’ and ‘Tribological properties’ sections. The variations of mechanical properties and in vitro behaviour of HA–CNT nanocomposites as a function of processing temperature, because of phase and microstructural changes, have been studied in detail by Xu et al. ²⁵¹ There have been controversial views in relation to the benefits of CNTs in biomaterial applications. On the one hand, CNTs were demonstrated to enhance cell proliferation. ^249,252 On the other hand, CNTs were found to exhibit toxicity in vivo. ^253,254 The toxicology effects of thermal sprayed CNT-reinforced biocoatings will be discussed further in the ‘Biocompatibility and cytotoxicity’ section.

Polymer-based nanocomposites

Polymer-based materials are widespread as protective coatings in harsh environments, low friction insulation coatings, and functional coatings because of their cost effectiveness, low density, chemical inertness, as well as high tensile strength. The inclusion of reinforcing filler results in composite coatings with enhanced mechanical and tribological thermal properties. ^255,256 However, conventional polymer processing techniques often require the use of volatile organic solvents and have limited on-site applicability for large surfaces.

The thermal spray process not only eliminates the need for a solvent, through the use of powder feedstock, but also offers the following advantages: (i) the capacity for on-site application and repair without the restriction of component size and coating thickness; (ii) the ability to operate under a wide range of environmental conditions; (iii) the ability to deposit polymers with high melt viscosities, e.g. high filler content polymer composites; and (iv) the capacity for the particles to spread out to form coatings without being fully molten, because of the high kinetic energy. A review on thermal spraying of polymers have been conducted by Petrovicova and Schadler. ²⁵⁷

Polymer-based nanocomposite coatings have been deposited predominantly via low temperature thermal spray processes; i.e. combustion thermal spray processes such as flame spray and HVOF and the cold spray process to limit polymer decomposition. Thermal spray polymer-based coatings have focused primarily on thermoplastics such as Nylon, ^258,259 polycarbonate, ²⁶⁰ and polyether ether ketone (PEEK), ²³¹ although thermoset polymer coatings such as epoxy ²⁶¹ have also been documented. Polymer fractions in such nanocomposite coatings vary from 5 to over 95 wt-%, ^86,261 with the remainder composed of reinforcing fillers such as silica, ^258,262 carbon black, ^258,262 nanodiamond ²⁵⁹ and α-iron. ²⁶¹

Silica or carbon black reinforced Nylon nanocomposite coatings deposited by flame spray and HVOF exhibited significantly enhanced crystallinity, which consequently increased mechanical and tribological performance of the coatings. Scratch and wear resistance of Nylon–SiO₂ nanocomposite coatings were improved by up to 35 and 67%, respectively. ²⁶² The increase of crystallinity was evidenced by the increase of melting enthalpy or heat of fusion; as measured by differential scanning calorimetry (DSC) and an increase of peak intensity and overall area under X-ray diffraction (XRD) spectra. ^86,263 Such improvement in the degree of crystallniity was postulated to arise from the initiation of heterogeneous crystallisation and changes in the polymer thermal characteristics by the reinforcement filler. ²⁵⁸

The crystallisation process did not cause any cracking or coating shrinkage because of the space restriction imposed by the nano-sized filler; i.e. small crystalline domains were concentrated in the vicinity of the nanoparticles. ²⁵⁸ Chemical modification of silica further improved the degree of crystallisation, where the degree of crystallinity increased in the order: hydrophilic SiO₂ < hydrophobic SiO₂ < carbon black. ^258,262 However, other contributing factors on the degree of crystallinity, such as degree of heating/melting of the polymer particles during in-flight and the substrate temperature profile, have not been taken into consideration and require further investigation.

Only two out of five crystal structures of Nylon were detected in the feedstock and coatings: (i) the thermodynamically stable triclinic α phase, and (ii) the metastable pseudohexagonal δ phase. ²⁶³ Pure Nylon and nanocomposite coatings did not retain the metastable δ phase upon deposition, whereas micrometre-scaled nanocomposite coatings preserved a low level of δ phase. Elimination of the δ phase was achieved by a relatively slow cooling rate after deposition, namely ∼50°C min⁻¹, thus allowing sufficient time for the δ phase to rearrange to the thermodynamically stable α phase. This effect indicated that the particles underwent some melting during deposition. Melting has been reported for Nylon–SiO₂ nanocomposite coatings, but no decomposition products have been detected. ^86,258,263

Deposition of nanodiamond reinforced Nylon coatings by HVOF indicated effective retainment of the nanodiamond phase after deposition, and did not result in any significant graphitisation of the nanodiamond. ²⁵⁹ The uniformity of nanodiamond distribution in the coatings was improved by purifying the nanodiamond through surface chemistry modification. The modification allowed conversion of surface functionalities to –COOH groups and, this was followed by treatment with dilute hydrochloric acid to remove metal impurities. ²⁵⁹ The surface –COOH groups form hydrogen bonds with nitrogen and oxygen atoms of the amide groups in the Nylon backbone chain. This reaction allows stronger interactions between the polymer and nanodiamond particles and a more uniform distribution of nanodiamond within the polymer matrix.

The viscoelastic behaviour of the polymer was modified by nanoparticle reinforcement, where the mobility of the polymer chains and conformation changes were limited by the nanoparticles. The confinement of the chains at the nanoscale level reduced the dissipation of energy during deformation processes. Hence coatings with high nanodiamond content tended to exhibit greater recovery at the end of an unloading cycle and smaller residual extension compared to pure polymer coatings. ²⁵⁹ The adhesion and peel strengths of the nanocomposite coatings were significantly improved by the polar bonding of SiO₂ ⁸⁶ and –COOH surface groups of nanodiamond ²⁵⁹ with the substrate.

Reports of the performance of other thermal sprayed polymeric nanocomposite coatings have been directed towards specific functional properties such as electromagnetic wave absorption for applications such as silent rooms and radar systems ²⁶¹ and anti-microbial properties for biomedical applications. ²³¹ Similar to mechanical properties, the extent of functional property enhancement of these coatings depend on attributes of the nano-components such as volume fraction, uniformity of distribution, particle size and morphology, and magnitude of the specific functional property. No polymeric decomposition products were reported in these functional coatings deposited by high velocity air fuel and cold spray. ^231,261

Mechanical properties

The four main criteria that are often used as measures of the mechanical performance of thermal sprayed coatings are (i) hardness, (ii) elastic modulus, (iii) fracture toughness, and (iv) bond strength. Nanocomposite coatings have generally shown marked improvements in all of these mechanical attributes compared to their conventional, microstructured composite coatings as well as unreinforced, single-phase nanostructured coatings. Nanostructured materials that were reported to benefit in these aspects include (i) carbide or oxide reinforced materials, such as Al₂O₃–TiO₂, Al₂O₃–ZrO₂, ZrO₂–Y₂O₃, WC–Co, B₄C–Al, and Cr₃C₂–NiCr, and (ii) carbon nanotube reinforced matrices, such as hydroxyapatite, Al₂O₃, and Al–Si. The mechanical properties of several nanocomposite coatings deposited by various thermal spray processes are summarised in Table 12.

A 20% increase in hardness has been reported in nanostructured Cr₃C₂–25wt-%(Ni20Cr) coatings deposited by HVOF compared to their conventional counterparts, ^265,266 while the hardness of (Al–Mg)–20wt-%B₄C nanocomposite coatings was almost double that of unreinforced nanostructured Al–Mg. ¹³⁴ Cold sprayed nanostructured WC–12wt-%Co. coatings exhibited hardness values between 11·91 and 20·13 GPa, ^135,137,200 which is comparable to that of sintered bulk WC–Co. ²⁶⁷ The increase of hardness values is often explained in terms of the reduction of grain size in the nanocomposites. ¹³⁴

However, bimodal hardness distributions have been observed for nanocomposite coatings deposited by conventional approaches, which exhibit bimodal microstructures, as discussed in the ‘Conventional approaches’ section. Distinctive hardness values were noted in the fully molten and partially molten (nanostructured) regions of the coatings, where the more porous partially molten zones exhibited a lower hardness value compared to the fully molten zones. ^120,121 Such behaviour indicates that the microstructure and density of the coatings have more important roles in determining the hardness values than the grain size when there are microstructural variations throughout the coatings.

The reported trend of elastic modulus values for nanocomposite coatings compared to conventional coatings has been inconsistent. While most nanostructured coatings were reported to have a decreased elastic moduli because of increased strain tolerance through irreversible deformation by particle sliding and microcracking, ^268,269 others were reported to have an increased modulus. ^183,270 Such discrepancies were attributed to the difference in the density and volume fractions of the nanozones in the coatings, as well as a sensitivity of nanostructured coatings to localised measurements and test parameters; e.g. load and indenter geometry.

Plasma sprayed nano-(ZrO₂–7wt-%Y₂O₃) coatings exhibited a more stable, but lower, elastic moduli when exposed to prolonged high temperature conditions compared to conventional coatings, although both coatings demonstrated a significant increase of elastic moduli at the beginning of heat treatment; ²⁶⁸ Fig. 35. The stabilisation of elastic modulus values was induced by the differential sintering behaviour of nanostructured coatings. The high sintering rate of the porous nanozones compared to the fully molten regions led to the opening of large micrometre-sized voids in the coatings, which counteracted the overall densification because of sintering. ¹¹⁹ Such porous nanocomposite coatings, which have low stiffness and are durable through thermal cycling, have been patented as thermal barrier and abradable coatings in applications such as turbine blades and the combustion chambers of aircraft and land-based gas turbines. ²⁷¹

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Figure 35.

Variations of elastic modulus values of nanostructured and conventional ZrO₂–7wt-%Y₂O₃ coatings when subjected to heat treatment at 1 400°C. Reproduced from Ref. 268 with permission from Springer Science and Business Media

The fracture toughness of plasma sprayed hydroxyapatite coatings was increased by 56% from 0·39±0·09 to 0·61±0·09 MPa m^1/2 when reinforced with 4 wt-% of MWCNTs. ²³⁴ The improvement of fracture toughness was attributed to the distribution and anchoring of CNTs to form bridge structures, which have excellent elastic recovery and could withstand high bending deformations. Similar effects were observed for MWCNT reinforcements in Al₂O₃. ²⁷² The CNTs form a web-like structure with the ceramic grains in the nanocomposites, which improve cohesion and inhibit grain growth by restricting grain boundary mobility. ²⁵¹ However, such cross-linking by CNTs could be damaged by post-deposition heat treatment above 1 200°C that consequently result in total loss or reduced reinforcing effects of CNTs.

Nanostructured Al₂O₃–13wt-%TiO₂ coatings also demonstrated a significant increase in fracture toughness as measured through crack growth resistance and interfacial toughness. ^124,173 While the long and well-defined splat boundaries in conventional coatings tend to provide an easy crack propagation path, the periodic disruptions of splat boundaries by the partially molten regions in nanocomposite coatings act to arrest or deflect crack propagation; thereby improving their fracture toughness. Similar crack bridging effects were observed through improvement in the interfacial toughness of Al₂O₃–13wt-%TiO₂ nanocomposite coatings. Improved toughness has also been attributed to enhanced interlamellar strength because of more homogeneous particle distributions. ¹⁶⁸

The bond strength of thermal spray coatings is governed by the cohesion strength between splats and adhesion between the coating and substrate. Determination of the cohesion and adhesion of Al–20wt-%B₄C coatings through shear testing indicated that the coating adhesion increases in the order: microstructured Al–B₄C, unreinforced aluminium alloy and nanostructured Al–B₄C. ¹³⁴ Reinforcement using micrometre-sized ceramic particles reduced the adhesive strength because the ceramic particles caused a decrease of surface area for matrix particle deformation at the coating and/or substrate interfaces. On the contrary, nano-sized B₄C ceramic particles were encapsulated within the aluminium alloy matrix, thus minimising the contact area of the ceramic particles at the interface and allowing effective bonding. The contact between splats was also improved in nanostructured coatings because of the refined splat size, resulting in the enhancement of cohesive strength. Similar bond strength improvements were noted through measurements by tensile adhesion tests in nanostructured ZrO₂–7wt-%Y₂O₃, Al₂O₃–13wt-%TiO₂ and WC–12wt-%Co coatings. ^123,137,179

Thermal sprayed durable and tough nanocomposite coatings have been implemented in industries such as aerospace, automotive, and marine. Nanostructured Al₂O₃–13wt-%TiO₂ coatings were approved by the United States Navy for shipboard and submarine applications such as arm weldments, bulkhead pivot arms, bearing sleds, front and aft door supports, magnet arms, socket and arm pivot pins, periscope guides, hydraulic pistons, and reduction gear sets. ^124,273 Thermal sprayed nanostructured coatings with enhanced mechanical properties, such as ZrO₂–7wt-%Y₂O₃ and (ZrO₂–7wt-%Y₂O₃)–Al₂O₃, are widely used as durable thermal barrier coatings for gas turbine engines in aircraft and power generation plants. ^274–276

Nanocomposite coatings with high hardness and fracture toughness incorporating 3–5 nm carbide or oxide grains embedded in an amorphous matrix have also been developed for aircraft components. ²⁷⁷ The coatings capitalise on the macroscopic ductility mechanism that is unique to nanocrystalline–amorphous composites through the large proportion of 1–2 nm nanograins in the amorphous matrix. Mechanically hard coatings containing dispersed nano-sized carbide particles, such as SiC, TiC, WC, or VC, have been patented as hard protective coatings for internal combustion engine components such as piston rings. ²⁷⁸ Hard nanocomposite coatings are also advantageous as wear resistant coatings, which will be discussed in the next section.

Tribological properties

There is a general trend reported in the literature that the wear resistance of a material is directly related to its hardness and toughness; i.e. wear resistance is increased when hardness and/or toughness increased. ^121,174,184 Such effects were evidenced in thermal sprayed nanocomposite coatings, where the enhanced wear resistance pertains to (i) the crack arrest and deflection by nanozones, ¹²⁴ and (ii) the enhanced interlamellar strength from homogeneous particle mixing. ¹⁶⁸ Depending on the wear mechanisms and the coating characteristics, the coating hardness and toughness influence the wear resistance to different extents. For instance, suspension plasma sprayed Al₂O₃–40wt-%ZrO₂ coatings exhibited a higher abrasion resistance compared to that deposited by HVOF owing to its superior hardness. ¹⁵⁸ However, the brittle nature of the plasma sprayed coatings has resulted in poor erosion resistance compared to HVOF coatings, which have high crack propagation resistance and toughness compared to plasma sprayed coatings.

Nanostructured ZrO₂–7wt-%Y₂O₃  =  coatings have shown a 20% reduction in wear rate compared to its conventional counterpart, although there has been no significant changes to the coefficient of friction. ²⁷⁰ In addition to the phase transformation toughening by stress-induced martensitic transformation of ZrO₂ from tetragonal to monoclinic, the enhanced wear resistance was also attributed to the narrowing of the defect size distribution, which led to the improvement of resistance to crack propagation. Incorporation of 4 wt-% CNTs reduced the wear volume loss of hydroxyapatite coatings by almost 50%. ²³⁵ The CNTs act as reinforcing and self-lubricating elements in the coatings, while improving the fracture toughness of the coatings through anchoring effects and energy absorption from protruding CNTs at the surface during abrasion.

The abrasion resistance of Al₂O₃–13wt-%TiO₂ coatings was improved eight-fold by using a nanostructured feedstock, compared to conventional coatings. ¹⁶⁷ The dense and finely dispersed nanozones, which were acquired through penetration of the semi-molten porous core by molten material during in-flight, contributed to the increased crack propagation resistance and wear resistance of the coatings. The wear mechanisms in conventional Al₂O₃–13wt-%TiO₂ coatings were identified to be grain dislodgement because of grain boundary fracture and lateral crack chipping, while wear in nanostructured coatings was only through grain dislodgement. ¹⁷⁰ Wear resistance of Al₂O₃–13wt-%TiO₂ coatings can be further improved by three-fold through the addition of CeO₂ and ZrO₂ particles; ^170,279 Fig. 36. Similarly, the addition of 50wt-% Fe₃O₄ as a solid lubricant in Al₂O₃–13wt-%TiO₂ coatings significantly reduced the coefficient of friction, which result in a four-fold increase in the sliding wear resistance and a three- to five-fold increase in the abrasive wear resistance. ²⁸⁰ Such self-lubricating wear resistant coatings find application in high pressure fuel pumps for direct injection engines that use cleaner, lubricant-free fuels.

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Figure 36.

Wear volumes of conventional and nanostructured plasma sprayed Al₂O₃–13wt-%TiO₂ coatings. Reproduced from Wang et al. ¹⁷⁰ with permission from Elsevier

On the contrary, the abrasive wear resistance of nanostructured WC–Co coatings was found to be inferior to that of conventional coatings. ^201,203 The wear rates of HVOF sprayed nanostructured WC–Co were between 1·4 and 3·1 times higher than coatings deposited using conventional powders. The deterioration of wear resistance was attributed to WC pullouts and decomposition. The wear mechanism of WC–Co coatings is through binder phase removal by abrasive particles. Therefore, larger carbide particles in conventional coatings are more resistant to wear because the binder phase must be removed to a greater depth before the particle can be removed. ²⁰³

On the other hand, the fine carbide particles in nanostructured coatings can be detached with the removal of only a small amount of binder, and the fine WC pullouts contribute to further abrasion. The nano-sized particles are also more susceptible to in-flight decomposition compared to conventional particles, which led to the decrease of the wear resistant WC phase and the formation of brittle amorphous binder phases in the coatings. ²⁰¹ Implementation of a multimodal feedstock consisting of a mixture of nano- and sub-micrometre-sized particles rather than a fully nanostructured feedstock improved the wear resistance of WC–Co coatings by over 30% compared to conventional coatings because of better retainment of WC particles. ^281–283 Such multimodal coatings have been patented as abrasion wear resistant coatings. ²⁸⁴

The erosive wear resistance of Al₂O₃–20wt-%(ZrO₂–8Y₂O₃) coatings was increased by about 50% through the addition of 10% nano-SiC particles. ¹⁸⁴ The SiC hard phase contributed to an increase of hardness and fracture toughness of the coatings, thereby improving the erosion resistance. Similar improvement in scratch resistance and sliding wear resistance was observed in polycarbonate and Nylon-11 coatings through SiC and/or carbon black nanoparticle reinforcement. ^{46,260,262,263} Coatings that incorporate modified hydrophobic ceramic reinforcing particles exhibited the best scratch and wear resistance because of enhanced polymer-ceramic adhesion and reduced agglomeration of the reinforcement particles. ^46,263 The scratch resistance increased as a function of decreasing reinforcement size ²⁶³ and increasing polymer particle size. ²⁶²

Hard and tough nanocomposite coatings can be used as wear resistant coatings for structural applications in a variety of industries. Wear resistant nanocomposite coatings can potentially increase component life up to five times. ^37,285 Self-lubricated, thermal sprayed nanocomposite coatings comprising a BN or Fe₃O₄ lubricant phase, a hard ceramic phase, and a ductile metal phase have been patented as low friction and wear resistant coatings for metal machinery. ^286,287 Thermal sprayed coatings containing nano-Cr₂O₃ and a secondary oxide or carbide immiscible phase are used as erosion and abrasion resistant coatings to enhance the durability of valve components. ²⁸⁸ Polymer coating systems consisting of solid lubricants such as high density polyethylene, polytetrafluoroethylene, graphite, or molybdenum disulphide, which are deposited by spray coating, dip coating, pulsed laser deposition, CVD, or electrochemical deposition, can potentially be adapted for thermal spray deposition and used as protective coatings on threaded joints for metallic pipes. ²⁸⁹

While dense nanocomposite coatings are widely used as wear resistant coatings, porous nanocomposite coatings with a friable structure are used as abradable coatings. Porous nanocomposite coatings are produced by retaining the porous nature of the nanostructured feedstock through incomplete infiltration of the non-molten particle core by the semi-molten region. ²⁷¹ Abradable coatings are typically porous composites that consist of a metal phase and a self-lubricating non-metal phase. Such coatings are particularly useful as abradable seals to reduce the gas path clearance in gas turbine engines and improve the efficiency and fuel consumption of aircraft and power plants. ⁷⁹ These coatings must be abradable yet mechanically stable to withstand high temperature operating conditions. Therefore, abradable coatings are often associated with low thermal conductivity thermal barrier coatings, which will be discussed in the next section.

Thermophysical properties

Heat conduction in solids occurs mainly through lattice waves; i.e. elastic or ultrasonic waves through phonon propagation. Heat conduction by lattice waves is influenced by the lattice defects, grain boundaries, and extended imperfections, which collectively determine the mean free path or attenuation length of the heat carriers. ²⁹⁰ However, at elevated temperature, the radiative component becomes a significant mode of heat conduction in addition to lattice waves. Heat transfer through radiation can be reduced by porosity or inclusions with a substantially different index of refraction than the radiating body. Therefore, the optimised microstructure for thermal conductivity reduction is a matrix of nanometre-sized grains with relatively large pores or other stable inclusions of about 0·5 μm diameter. ²⁹⁰

Comparing nanostructured thermal spray coatings to their conventional counterparts, nanostructured coatings exhibit a lower thermal conductivity and diffusivity because of their small grain size and homogeneous pore distributions. ^291,292 The attenuation length decreased when the grain size decreased, hence effectively reducing grain boundary scattering and thermal conductivity; ²⁹⁰ Fig. 37a . The smaller and more homogeneously distributed pores and thinner splats in nanostructured coatings provided a larger boundary volume for effective phonon scattering and stronger interfacial thermal resistance, which consequently reduced the thermal conductivity. ²⁹² The equiaxed grains in nanostructured coatings with tortuous zigzag-like boundaries also contributed to a greater thermal resistance compared to conventional coatings, which exhibit columnar grains that are aligned parallel to the direction of heat flow. ²⁹²

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Figure 37.

a Theoretical thermal conductivity of ZrO₂–7wt-%Y₂O₃ coatings as a function of temperature and grain diameter (reproduced from Klemens and Gell ²⁹⁰ with permission from Elsevier) and b evolution of thermal conductivity values of ZrO₂–7wt-%Y₂O₃ coatings during sintering at 1 400°C (reproduced from Lima and Marple ²⁶⁸ with permission from Springer Science and Business Media)

The thermal conductivity of plasma sprayed conventional ZrO₂–7wt-%Y₂O₃ coatings was reported to be almost twice that of nanostructured coatings. ²⁶⁸ A similar reduction of thermal diffusivity was observed for nanostructured ZrO₂–7wt-%Y₂O₃, ^119,291 and Al₂O₃–50wt-%ZrO₂ coatings, ²⁹³ indicating that the coating density is not the determining factor for the decrease in thermal conductivity. Such effects are the result of effective phonon scattering and interfacial thermal resistance associated with the decrease of grain size and splat thickness. The influence of interfacial thermal resistance on thermal conductivity is evidenced by the decrease of thermal conductivity through deposition of thin, multi-layer coatings. ^294,295 The thermal expansion coefficient of nanostructured ZrO₂–7wt-%Y₂O₃ coatings was also found to be higher than that of conventional coatings. ²⁹¹

Exposure to high temperature resulted in a substantial increase in thermal conductivity for both conventional and nanostructured coatings. ^119,268 The increase in thermal conductivity is attributed to sintering effects, where the pores were healed and the contact between splats improved. However, the thermal conductivity growth rate for conventional coatings was almost five times higher than that of nanostructured coatings; ²⁶⁸ Fig. 37b . Such differences in the thermal conductivity growth rate are attributed to the differential sintering behaviour of nanostructured coatings as discussed in the ‘Mechanical properties’ section. The stable thermal conductivity of nanostructured coatings at elevated temperature places them as a viable candidate for high temperature applications.

Suspension plasma sprayed Al₂O₃–40wt-%ZrO₂ coatings exhibited a higher thermal conductivity compared to coatings deposited by plasma spraying of agglomerated feedstock. ¹⁵⁸ The high thermal conductivity was attributed to low interfacial thermal resistance because of the more intimate splat contact and low porosity in SPS coatings. The improved bonding, in addition to increased strain tolerance, contributed to a longer thermal cycling life of SPS and SPPS coatings. ^179,181 On the other hand, the coatings deposited by suspension HVOF spray exhibited a lower thermal conductivity than its plasma sprayed counterpart because of a high amorphous phase content and limited particle bonding because of poor melting in the low temperature flame. ¹⁵⁸

The thermal cycling life of nanostructured ZrO₂–7wt-%Y₂O₃ coatings was 2–4 times higher than that of conventional coatings. ^296,297 Nanostructured coatings exhibited significantly different thermal shock behaviour than conventional coatings, where crack propagation in nanostructured coatings occurred through intergranular fracture while that in a conventional coating was through intergranular and transgranular fracture. ²⁹⁶ The increase of crack path tortuosity through intergranular fracture enhanced the thermal shock resistance of nanostructured coatings. During thermal cycling of nanostructured coatings, vertical cracks propagated from the surface towards the bond coat without any horizontal propagation at the coating interface because of the presence of nanozones. ²⁹⁶ Such vertical cracks were not found in conventional coatings and were replaced instead by horizontal cracks near the coating interface.

The use of low thermal conductivity nanocomposite coatings as advanced thermal barrier coatings can potentially improve the thrust and lower the fuel consumption of gas turbine engines. The thermal barrier coatings provide an additional layer of insulation to engine components such as turbine vanes and combustors. An example of such a nanocomposite coating is given in Ref. 298. A thermal barrier system that consists of two layers of ZrO₂–7wt-%Y₂O₃ coating with a nanostructured inner layer was claimed to improve the ductility of thermal barrier coatings. ²⁷⁵ Nanocomposite coatings with high thermal stability can also be used as abradable coatings to minimise clearance in high temperature applications.

The temperature distribution in a thermal barrier system is depicted in Fig. 38. Such nanocomposite coatings essentially exhibit vertical cracks, ²⁷⁶ and/or a porous nanostructure, ²⁷¹ which maintains the low conductivity and stabilises the thermal diffusivity after several hours of thermal exposure. An intermediate coating between a bond coat and a ceramic top coat, which consists of NiCoCrAlY and Al₂O₃ nanoparticles, has been implemented to reduce thermal stresses that arose from different thermal expansion coefficients between the bond coat and top coat. ²⁹⁹

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Figure 38.

Schematic of temperature distribution across thermal barrier system

Electrical properties

The electrical properties of nanocomposites have been widely capitalised, either for their conductivity ³⁰⁰ or insulation, ¹⁹ in applications such as electronics, ³¹ energy conversion and storage, ³³ and tissue engineering. ³⁰¹ However, reports on the electrical properties of thermal sprayed nanocomposite coatings have been scarce. The literature for these coatings has been more focused on the coating microstructure rather than presenting evidence of their electrical properties.

Suspension plasma spray has been presented as a potential solution for producing a bilayer anode–electrode nanocomposite system for SOFC applications. ³⁰² The bilayer system consists of a porous Ni–(ZrO₂–8wt-%Y₂O₃) layer for the anode and a dense ZrO₂–8wt-%Y₂O₃ layer as the electrolyte. The different microstructures were attained by manipulating the in-flight droplet's surface tension and viscosity through the use of different solvents; i.e. water suspensions to produce a dense microstructure and alcoholic suspensions for the porous microstructure. A functionally graded TiO₂–hydroxyapatite nanocomposite coating has also been deposited by SPS to improve the interfacial bond between hydroxyapatite and the implant material. ^238,239 The electrical conductivity of TiO₂, which has been measured in bulk materials, ³⁰³ and nanostructured coatings, ^304–306 could potentially be used to direct cell growth in biomedical implants. ³⁰¹ However, this approach has yet to be verified because of a paucity of experimental data.

Nanostructured dielectric coatings for applications in low frequency electrical components and as insulation layers in embedded three-dimensional circuitry have also been developed via thermal spray. ³⁰⁷ Nanocomposite Al₂O₃–13wt-%TiO₂ and Al₂O₃–5vol.-%Ni coatings were found to have increased electrical permittivity compared to pure nano-Al₂O₃ coatings; while the incorporation of nano-NiO, SiO₂ and glass were found to have decreased the permittivity. The varying effects of the different materials on permittivity were related to the presence of porosity, extent of α-Al₂O₃ phase retainment and the surface roughness of the coatings. Conductive nanocomposite coatings have also been developed in the same study, where the incorporation of 5 wt-% of nano-WC doubled the electrical conductivity of pure copper because of the suppression of oxide formation. ³⁰⁷

Similarly, the electrical conductivity of silicon carbide, SiC, coatings, which are widely used as functional coatings in electronic and opto-electronic devices, was improved by the incorporation of nano-sized titanium nitride, TiN. ¹³¹ Silicon carbide–titanium nitride nanocomposite coatings deposited by cold spray were composed of a partially agglomerated SiC insulating phase because of their high surface area and surface energy. The conductive TiN phase in this nanocomposite coating was homogeneously distributed to form a net-like structure that contributed to the electron transport properties of the coatings. ¹³¹ The electrical resistivity of the coatings was decreased with the increase of conductive TiN contents.

A nanocomposite system integrating a photoelectric conversion function and electron storage function into a single device, which could potentially be used for photo-cathodic corrosion protection and solar energy conversion and storage, was produced by thermal spraying TiO₂–50wt-%Fe₂O₃ coatings. ³⁰⁸ The reduction of TiO₂ particle size to nanoscale has enhanced the photocurrent of the coatings by more than 10 times. Such enhancement was attributed to the increase of the particle surface area and quantum size effect. On the other hand, the use of nano-Fe₂O₃ doubled the charge and discharge capacity of the coatings, which contributed to the improvement of the electron storage function. The enhancement of electron storage capacity was associated with the increase in reaction rate from the larger surface area of nano-Fe₂O₃.

Plasma spray was used to create metastable ceramic phases for SOFC applications through rapid quenching. ³⁰⁹ The fuel cell components consist of three nanostructured plasma sprayed layers: (i) nanoporous lanthanum strontium magnesium, La_1−xSr_xMnO₃, as a nanoceramic cathode; (ii) dense metastable YSZ, ZrO₂–8wt-%Y₂O₃, as the electrolyte; and (iii) porous Cu/Ni–(ZrO₂–8wt-%Y₂O₃) nanocomposite as the anode. Porosity was introduced into the anode and cathode layers to enhance the catalytic activities at the electrodes. The plasma sprayed nanocomposite SOFC reduced the operating temperature from 800 to 600°C by increasing the rate of electrochemical reactions through enhanced catalytic activity and faster transport of negative ions. Nano-gas channels and porosity were introduced to the SOFC's nanocomposite anode by plasma spray to further enhance the electrochemical activity and retard nickel aggregation under the high temperature environment. ³¹⁰ The conductivity of Ni–(ZrO₂–8wt-%Y₂O₃) nanocomposite manufactured by plasma spray compared to its conventional microstructured counterpart is depicted in Fig. 39. Nanostructured samarium doped ceria deposited by SPS was also proposed as an alternative SOFC electrolyte material for a low operating temperature environment. ²⁷⁴

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Figure 39.

Conductivity of (ZrO₂–8wt-%Y₂O₃)–Ni nanocomposite anode as a function of temperature, compared to that of conventional microstructured anode fabricated by tape casting ³¹⁰

Nanocomposite semiconductor materials were also developed by thermal spray for opto-electronic applications. ³¹¹ The coating is composed of a polymorphic matrix, such as silicon or germanium doped derivatives, with dispersed metastable nanocrystals formed during the thermal spray process. Nano-fillers such as CNTs are also known to produce composites with enhanced electrical properties. ^19,31,215 Despite many studies on thermal sprayed carbon nanotube composites, the potential of these composites in relation to their electrical properties is yet to be explored.

Other properties

Feedstocks

The thermal spray process can be segmented into five distinct groupings as indicated in Fig. 40. The presented sequence from 1 to 5 is intended to be logical; i.e. the thermal spray process starts with an appropriate feedstock and finishes with testing and evaluation for quality control purposes before being implemented towards specific applications as indicated in box 6.

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Figure 40.

Input criteria that determine the outcomes for nanoscale materials manufactured by thermal spray methods

Nanostructured coatings exist in some microstructural form for the majority of thermal spray coatings perforce of the rapid solidification process. ^52,338 Thus, microstructural artefacts; e.g. splat sizes, porosity, cracks, phase distributions and grain sizes are likely to form. However, the future outlook demands such features to be dominant and manufactured in a controlled fashion rather than be a curious and scientifically interesting anomaly. The most appropriate feedstock is consolidated from nano-particulates that are produced via a chemical route. The art in the thermal spray processing is to adjust the spray parameters so that only partial melting arises in order to maintain nanostructures that are consolidated into an integral coating by partial melted particles. ^75,79,103

The SPS and SPPS methods show considerable promise in lowering the feedstock cost barrier and achieving nanostructured deposits; however, there are associated challenges with regard to attaining high deposition efficiencies and high rates of deposition; ^104,105,339 characteristics that are both vital for manufacturing environments. It is foreseen that specific applications, possibly in the thermal barrier coatings or electronic devices sectors, will be primary drivers to grow these methods.

The principal challenge in the area of feedstock is to attain nanostructured deposits, therefore, future developments will, most probably, involve a combination of conventional and solution/suspension materials. Thus, it is suggested that (i) duel injection port strategies and (ii) use of a single injection port, but slurry feedstock consisting of traditional solid and low viscosity mixtures, will pave the way for increased production routes.

Associated with these industrial needs are those of (i) an ability to scale up feedstock production and ensuring their long storage capability, especially if any fluid components are present, and (ii) safe guards and controls with respect to occupational health and safety. The issue of cost will always be present and scale-up in specific niche areas, which are suggested above, will drive cost down.

Processes and manufacturing

The traditional and recent thermal spray processes in box 2 of Fig. 40 have been deployed successfully to manufacture nanostructured materials. The main issue, though, boils down to manufacturing constraints that arise under these well-known processes. Thus, these processes are line-of-sight, exhibit relatively low deposition and spray efficiencies, and are inflexible with regard to the spray footprint that controls the deposit resolution. ⁵² These factors are all adverse to the manufacturer economics that drives innovation.

The promise of nanoscale processing via thermal spray is that there may be an opportunity to break down the barriers, also described as ‘challenges’, posed in Fig. 40; especially by employing multiple processes. For example, sol-gel infiltration ³⁴⁰ of an existing conventional thermal spray coating via dipping or simple aerosol spray methods that is followed by a post-process such as (i) arc plasma heating (with or without the thermal spray device), ³⁴¹ (ii) conventional furnace heat treatment, ^342,343 or (iii) laser glazing. ³⁴⁴ As well, there is a large market for near-net shape manufacturing processes; therefore, additive manufacturing by means of thermal spray deposition is a distinct opportunity.

It is recognised that there is resistance to changes in any engineering practice. Thus, the technical and financial advantages need to be articulated clearly before the adoption of such practices.

The microstructure

The microstructure is driven by the need for certain intrinsic and extrinsic needs. The prior knowledge for thermal spray coatings indicates that the physical characteristics are determined by the extrinsic characteristics. That is, the physical properties depend strongly on the defective microstructure; ³⁴⁵ which accounts for their highly variable nature.

The opportunity to manufacture nanostructured coatings offers potential to enhance physical attributes such as modulus, ductility, thermal conductivity, and magnetic and electrical properties. Such properties may be tailored by adjusting the microstructure. For instance some strategies include (i) modifying the feedstock with regard to chemistry and intrinsic size, ³⁴⁶ (ii) varying the thermal spray process so that the feedstock experiences appropriate thermal and velocity conditions to achieve the desired microstructure, ¹⁹⁹ and (iii) using other manufacturing techniques, such as sol-gel infiltration, ³⁴⁰ laser post-processing, ²⁰³ and sintering techniques, ^268,341 to modify the as-received thermal spray coating.

The overall notion is to re-design the engineered coating on the basis of the desirable microstructure, rather than being constrained by conventional feedstocks and processes. Therefore, the microstructure is transformed into being the prime design criteria and not a fortuitous accident that derives from the manufacturing process.

Testing and evaluation

Testing of nanostructured materials is well-developed. For example, mechanical properties may be determined by nano-indentation ^347,348 and scratch techniques; ^184,263 phase structure by means of XRD, ^122,123 Raman spectroscopy, ³²⁶ transmission and scanning electron microscopy; ^217,237 and roughness via atomic force microscopy. ^349,350 These methods can be classified as ‘conventional and routine’ techniques.

The nature of all test methodologies on thermal spray deposits is contingent on the understanding that the property measure is extrinsic in nature. Therefore, the measured values are highly variable because of the inhomogeneous structure of these coatings. The exceptions to this understanding would incorporate processes that create dense coatings because these lamellar microstructures are typically disrupted. The extrinsic nature of coatings implies that properties, atypical of the original materials, can be realised and herein lies technical opportunities. Thus, the design engineer is no longer constrained by text book values for material properties because these are regulated by the thermal spray nanostructure. Furthermore, the statistical variability of the nanostructure will be related to the phase distribution; thereby leading to known property distributions.

Applications in the future

The current (2014) global thermal spray industry for the domestic market is valued at about 7·1 B. ³⁵¹ The market for these traditional applications of thermal spray coatings include ^52,53 (i) transportation systems such as aerospace, automotive, naval and railway networks; (ii) heavy industry such as remanufacturing for dimension restoration, and rolls in the steel, papermaking, and printing industries; (iii) prostheses for the bioengineering market; and (iv) the energy producing sectors such as land-based turbines, geothermal sources, wind energy, and solid oxide fuel cells. These markets will expand with the use of appropriate nanoscaled structures because of the improved performance gains that evolve from the nanostructure.

The more critical growth markets for the future are those that will be uncovered because of the new materials and thermal spray processes that will be discovered. Extension of cold spray methods to the application of nanostructured ceramics and cermet composites would open up opportunities for advanced applications where wear control is an issue. In an alike fashion, SPS and SPPS methods are on the cusp of being of industrial importance and any shift to creating nanostructures will accelerate their commercial adoption.

Boxes 6–9 allude to the need to create new market sectors that are outside the traditional applications for thermal spray deposits. It is especially important to identify future markets such as (i) sensor arrays, (ii) opto-electronic devices, (iii) nanostructured coatings for energy harvesting, (iv) thin coatings as masking materials for thin film technologies such as PVD and CVD, and (v) nanostructured biomimetic materials. These markets are currently being explored with regard to thermal spray technologies and bode well for a healthy future.