Effect of post-processing treatments on mechanical performance of cold spray coating – an overview

2.1 HT

Huang et al. [15] investigated the impact of cold sprayed coatings and heat-treated densities (aluminium and copper) and a porous coating of titanium and SS316. It was found that greater dislocation density and a few micro- and nano-level particles were found in the coated sample due to the massive velocity effect and extreme plastic deformation. On the other hand, the porosity percentage has decreased significantly at lower heat-treated temperatures for denser coatings [15]. This is mainly owing to the phenomenon of diffusion, through which the inter-particle interfaces and small pores are closed [13]. Diffusion was not quite as intense in porous as compared to dense coatings, since there was debilitated space between the bond particles, so there is no reduction in porosity (Figure 1(b)) due to residual oxides and natural oxide film on the surface of the particles [13]. Even at the highest rise of the HT temperature, a higher amount of diffusion takes place, which results in the removal of the splat boundary and even grain formation in the coatings owing to recrystallization as shown in Figure 1(c). Furthermore, as the HT temperature rises, the grain in the coatings expands and the flaws in the as-sprayed coating shrink, as shown in Figure 1(d). Figure 1 illustrates the schematical microstructural view of CS dense and porous coatings under heat-treated condition. From the figure it is evident that mechanical characteristics of deposit were improved for dense coating with a reduction of porosity in coating structure compared to porous coating. It is investigated and acknowledged by Huang et al. [15].

Figure 1

Schematical microstructural view of CS dense and porous coatings under heat-treated condition: (a) powder coating; (b) diffusion phase; (c) recrystallization phase; (d) grain growth phase [15].

Furthermore, the dislocation density of the deposit in the substrate was reduced with enlarged grain size owing to the incidence of restoration on the microstructure, which was demonstrated through electron backscatter diffraction (EBSD) and transmission electron microscope (TEM) examination. Figure 2 illustrates that heat-treated Inconel 718 coatings show lower microhardness compared to as-sprayed coatings. This is due to a reduction in the dislocation density as it is illustrated under Kernel average misorientation (KAM) [18]. On the other hand, as-sprayed Inconel 718 delivers greater dislocation density, and straight bands are formed under the damaged splats because they are plastically deformed and analysed through TEM. Then the sample is heat treated. Dislocation is re-uniformed into the cell structure and produces few grains of recrystallization. Aside from this, grains with higher band thickness were present in the heat-treated Inconel 718 substrate. As a conclusion, the dislocation healing, growth of grains, and stress were eliminated with a reduction in the hardness. Further, the samples were in heat-treated conditions [17,20]. Figure 2 examines the mechanical properties of the Inconel 718 coatings. It delivers lesser hardness after the heat-treated process with various CS processing gases such as nitrogen and helium. Yin et al. [13] investigated the effect of SS316L coating on Al alloy substrate. The thermal treatment was conducted on two atmospheric conditions of air and vacuum. They revealed that the air annealing does improve the tensile strength of the SS16L coating surface compared with the vacuum environment. During air annealing treatment, new phases were developed at the interface region, mainly consisting of Cr, Mn, O, and Si. The probable phases are Mn₂CrO₄, Cr₂O₃, and SiO₂, which are usually of 316L steel exposed to air. When oxygen can penetrate through the inner-particle boundary regions, it can induce the formation of oxides at inner region of the coatings, which can act as an impediment to the formation of inner-particle bonds, resulting in lower cohesion behaviour of deposit and poor mechanical properties. But at the vacuum environment interface region oxidation does not take place, it leads to improvement of coating characteristics. The same behaviour was obtained by Meng et al. [17] on the SS304 coatings in a vacuum environment. After HT, the poorly interfaced particles were cured. Lee et al. [23] and Wang et al. [24] revealed that high temperature HT on CS coating shows higher potential to produce intermetallic products due to atomic propagation of various compounds and recrystallization. Further, particles such as nickel or titanium or iron are added to aluminium matrix. It is heat treated and produces Al–Ni, Al–Ti, or Fe–Al elements. Spencer and Zhang [10] investigated that the elements produced within the deposits, a layer of aluminium with magnesium elements, were produced on the interface of the aluminium deposit while coated on the Mg substrate after HT. As a result, the mechanical properties of the aluminium and magnesium interlayer are the same as those of magnesium alloy, as well as the electrochemical behaviour. Nonetheless, Kirkendall holes were observed on the aluminium with magnesium matrix coatings after successful HT. Due to the faster diffusion of aluminium in the intermetallic layer compared to magnesium, the vacancy will be travelled in the opposite direction due to the influence of diffusive flux, and the vacancy will be filled through the aluminium deposit and the aluminium–magnesium product interface. Yu et al. [14] examined the effects of HT on the various atmospheres for an hour. After HT, the vacuum environment improved harder compared to the other environments. This is due to better diffusion taking place between the particles. The same behaviour was obtained by Ren et al. [31] with reduced porosity after the HT process. Also, Tang et al. [22] investigated the Ti–WC deposit. They revealed that the pores in the deposit were reduced after HT. This improved the mechanical and tribological behaviour of the deposit. This occurred as a result of the TiC phase’s development. Sun et al. [21] studied Inconel 718 coatings at 990°C in a vacuum environment for 10 min using a furnace and eddy current HT. They revealed that the process is excellent at improving atomic diffusion and mass transformation along the deposit. It leads to an improvement in the physical properties of the deposits. Wong et al. [44] also investigated Inconel 718 coatings at 950–1,250°C for 1–2 h in an environment containing 10% hydrogen gas and 90% argon. The tensile properties of the deposit were improved by around 63% after HT when compared to the BM. Li et al. [41] investigated the Cu–4Cr–2Nb coatings in the vacuum environment for 2 h. They revealed that the unique Cr₂Nb phases were developed after HT. It will improve the pinning impact of grain growth and dislocation motion of the particles. This enhanced the hardness of the coating. Also, Yang et al. [24] and Lee et al. [23] obtained a similar intermetallic layer. This intermetallic layer will strengthen the coating integrity. This results in the improved mechanical behaviour of the coatings [26,27,32].

Figure 2

(a–d) KAM outcomes for Inconel 718 CS coatings with various processing gases with and without heated-treated samples acquired through EBSD analysis: (a) CS through nitrogen, (b) CS through helium, (c) HT made on CS through nitrogen as a processing gas, (d) HT made on CS through helium as a processing gas, and (e and f) the microhardness of the above CS and heat-treated Inconel 718 deposits [18].

Yu et al. [62] investigated HT on copper deposits. The vast majority of weakly bound interfaces have been retrieved by metallic diffusion, and the deposit holes have started to take on a smoother and circular morphology. Figure 3 illustrates that it was discovered using the EBSD technique. The splat to splat boundaries in the coated sample vanished, and the extremely fine nanocrystalline structure was changed to the well-tiny equiaxed grains. This phenomenon is confirmed by various studies [40–45]. After HT, the pores in the deposit become circular in shape. That will minimize the concentration of stress along the pores and cracks of the deposit. Mechanical characteristics analysis and failure surface evaluation are shown in the figure. It demonstrated that, following higher temperature annealed samples, the plastic behaviour of the CS deposit is nearly identical to that of the bulk metallic sample. According to a fracture morphological examination, the deposit dimples after the annealing process shrank compared to the cold sprayed deposit. As was mentioned previously, annealing significantly alters the CS deposit tensile strength and ductile behaviour of the deposit. Figure 4 shows the microhardness of the deposit at various HT temperatures. As shown, the microhardness value reduces as the temperature rises, particularly at higher values. The four annealing steps are associated with the modification of microhardness. Typically, at lower temperatures, even just a portion of the work hardening is lowered, and the modifications to the grains and grain boundary are minimal. At higher temperatures, HT causes equiaxed crystals to form in the deposit and to develop when the annealing temperature and duration are increased. This significantly decreases the deposit hardness [62,63].

Figure 3

EBSD-IPF morphology [62]. (a) CSed copper deposit and (b) after heat-treated copper deposit.

Figure 4

Hardness behaviour of the heat-treated CSed A380 coatings [63].

However, many of the investigations into the HT techniques show that this is not enough for CS overhaul in industrial applications. Induction HT shows better results. In summary, HT is used to improve atomic diffusion between particles–particles and the coating–base substrate. It leads to an improvement in the coating’s ductile nature with greater improvement in the cohesive and adhesive behaviour of the deposit. But HT will change the microstructure of the coatings with the growth of grains, and the phase transition in the coatings is unavoidable. The advantages of using this HT is to improve the mechanical properties of samples for the light alloy material. However, various studies on heat-treated post-treated samples are reported as illustrated in Table 1.

2.2 FSP

Huang et al. [27] used EBSD to investigate the microstructural behaviour of CS and FSP-treated Cu–Zn coatings, as shown in Figure 5. The properties of the coatings are illustrated by the higher angles of 21.5% and lesser than the angle of 77.5% of the grain boundary. This is due to greater dislocation of the density with intense grain deformation. Then it is post-treated through FSP, which produces a higher angle of 90.6% and a lower angle of 9.5% of the grain boundary, owing to recrystallization with the plastically deformed condition at the time of FSP. Thus, the improvement in the high angle of grain boundaries after FSP results in an enhancement in the mechanical, wear, and corrosion properties. Furthermore, Huang et al. [27] examined phase transition of copper–zinc deposit under the CS process. They revealed higher alpha phases (red) in the deposit compared to beta phases, as seen in Figure 6. Moreover, the beta and gamma phases are formed after FSP because of the thermomechanical impact. The gamma phases produced on the surface are the result of the poor coating characteristics after FSP. As a result, caution must be exercised when applying FSP to CS deposit builds of thermally sensitive materials. Peat et al. [46,49] investigated on WC–CoCr/Al₂O₃ coatings. They revealed that after FSP, the reinforced particles uniformly scattered into the metal matrix composite coatings due to the shear force developed by the FSP process. At the time of FSP, the deformation occurs more in the agglomerates, which results in the scatter of carbide particles on the matrix phase. Finally, after FSP, the CS of metal matrix composite coatings shows the various conditions of refinement throughout the deposit. Since excellent refinement takes place on the top surface due to shear force produced by the tool and the middle and bottom surfaces of the deposit, the condition of refinement is lower owing to the reduction of shear force [46]. The dispersion of the as-coated WC–CoCr agglomerates and refinement of the ceramic particles. The FSP enhanced the hardest homogeneity and also the anti-erosion characteristics of the deposit [49]. Huang et al. [52,53] investigated the mechanical behaviour of a friction stir processed SiC-reinforced metal matrix composite deposit. The mechanical properties were lower after FSP compared to cold sprayed coatings owing to SiC particles being broken into various sizes. From this, while adapting the FSP, the process parameters need to be optimized properly to get the excellent microstructure and enhanced properties.

Figure 5

(a and b) EBSD outcomes: (a) cold sprayed, (b) friction stir processed deposit with black, white, and yellow lines are corresponding to high angle (θ ≥ 15°), low angle (15° > θ > 2°) grain boundaries. (c and d) Misorientation angle distributions: (c) CS, (d) FSP deposit [27].

Figure 6

BSD phase outcomes of (a) CS and (b) FSP Cu–Zn alloy coatings [27]. Here red represents α phase (Cu₃Zn), yellow represents β phase (CuZn), and blue represents γ phase (Cu₅Zn₈).

Table 2 illustrates the outcome of post-treated CS coatings. In FSP, the high heat input during the process produces very high dissemination of intermetallic phases, which is noticed in the microstructure. The same phenomena were observed in Ni and Ti particles as a result of development of Ni–Ti intermetallic phases [28]. Khodabakhshi et al.’s [29] investigation of the cold sprayed AA7075 coating on the AZ31B mg alloy after it is subjected to the post-process treatment (FSP) shows that the interface between coating and substrate shows the diffusion development of materials between the deposit and BMs with a small layer of intermetallic phase developed under friction formed through heat and plastic deformation. It provides the material intermixing. They observed that after FSP, a similar intermetallic layer was developed on the titanium coating with (Al₃Ti) layer and Ni50–Ti50 coating with (Ni–Ti) layer by Khodabakhshi et al. [50] and Huang et al. [28]. This will improve the wear resistance, mechanical, adhesion, and cohesion strength of the coatings [28,29,50]. Yang et al. [51,47] examined the AA2024/Al₂O₃ coatings. After FSP, the alumina particles were redistributed around the matrix phase and the alumina particles’ size were reduced. This improved the interface bonding between the matrix and the ceramic phase. The same phenomenon was obtained by Hodder et al. [48]. This leads to improve the hardness, tensile strength, and corrosion resistance of the coatings.

Huang et al. [27] illustrated that the FSPed sample structure is entirely denser in comparison to the CS deposit, as seen in Figure 7. The FSPed sample has eliminated all porosity nanoholes, fissures, and weaker binding interfaces that existed in the as-sprayed deposit. In an FSPed specimen, the native in-splat element crystal grains are transformed into well-fined equiaxed crystalline. Ashokkumar et al. [64] have mentioned that FSP is a successful technique for low-pressure cold sprayed AA2024/alumina metal matrix composite deposit on AZ31B Mg alloy. The CSed coating has a microhardness of 135 HV after FSP. The microhardness of the coatings was highly improved to 155 HV as shown in Figure 8 [64]. The agglomerates and initial bonding will be broken by the shear force caused by the FSP tool. The deposit will possess an even distribution of the alumina elements. The distribution of alumina particles in the FSPed coating will be more homogeneous and the particles will be finer than the as-sprayed coatings. Also, the corrosion resistance of the FSPed sample shows excellent resistance as shown in the figure. This is due to the enhanced surface condition of the FSP [49]. After repeated FSP operations, the as-sprayed coated microscopical hardness was measured using nanoindentation, and a hardness mapping of the region under examination was created, as seen in Figure 9. The figure demonstrates how it differs from the as-sprayed deposit. The homogeneity is highlighted by the FSPed deposit hardness. The tribological and erosional behaviour [28] of the FSPed sample is enhanced by its increased hardness.

Figure 7

Optical microscopical view of the cold sprayed and FSP [27]: (a) as-sprayed deposit and (b) friction-stir processed deposit.

Figure 8

Hardness characteristics of the cold sprayed and FSP [64] indention morphology: (a) as-coated sample, (b) FSP-treated sample, and (c) microhardness ranges of coated and FSP-treated sample.

Figure 9

Hardness mapping of the examined region [49]: (a) cold-sprayed WC–CoCr-reinforced metal matrix composite coating and (b) after FSP.

In summary, FSP is a better post-treatment to modify the microstructure of the CS deposit and enhance the bonding between the particles. The unique feature of the FSP is to improve the interface bonding between the matrix and the secondary phase of the deposit. It resulted in an improvement in the mechanical, electrochemical, and tribological properties.

2.3 SP treatment

In SP, the particles in the coated sample further get deformed, and the surface roughness of the deposit is lower due to SP treatment. Table 3 illustrates the cold sprayed coatings after SP treatment with various deposits. Moridi et al. [32] investigated the CS Al 6082 deposit and found a high ratio of nanograins in the coating structure after the SP [54] because of intense plastic deformation, and also very fine grains (100–200 nm) were observed in between the particle–particle interface of the cold sprayed Ni deposit as shown in Figure 10, with micron-size grains developed in the centre part of the coated area after post-treatment of SP. The micron size is modified into nanograins. As a result, after SP, Ni coatings have improved the electrochemical and hardness properties [33–37]. Ghelichi et al. [30] investigated the influence of post-treatment on CS coating after SP techniques. The findings indicate that after being cold sprayed, the SP was not effectively developed and had a significant impact on residual stress. SP just shows a limited impact on minimizing surface residual stress with a rise in the intense compression stress as shown in Figure 11. Since the CS deposit is not perfectly continuous, it forms coatings with higher flaws like particle and the bonding is poor. However, by ensuring peening, a significant portion of kinetic energy is expended on destroying the coating and removing some portions of the particles instead of causing work hardening. The particle removal may lead to the relaxation of residual stress. Moridi et al. [32] recommended that the elimination of deposition particles from the coatings be performed through the customized parameters. It is necessary to work the SP process with gradient parameters initially carried out at fewer process parameters with improvement in the coating density and interparticle bonding strength. Furthermore, the parameters of the process are increased, which results in the development of high compression stress with the formation of nanograins in the coatings.

Figure 10

EBSD mappings of the cross-sectional view of cold sprayed Ni coating: (a) Euler angle mapping and (b) pattern quality mapping of the same region [33].

Figure 11

Stress distribution curves of (a) cold sprayed and (b) SP process [30].

Lu et al. [65] investigated a CSed aluminium coating on the LA43M Mg alloy substrate. According to the analysis, the SP deposit surface morphology has been vastly enhanced over that of the CSed deposit, as seen in Figure 12. After SP, the porous structure in the deposit has decreased to an extremely low amount. As a result, there is a firmer bond between the particle and another particle and between the particle and the BM. The aluminium elements that were spherical at first vanished, and the flatness of the deposit was greatly enhanced. As the peening particles have a considerably larger diameter than the spraying particles, they will impact the coated particles and bounce. The CS deposit can be totally dense following the SP process. The SP particles and coating materials are blended in various fractions. The findings demonstrate that the porosity of the CS deposit diminishes as the SP particle fraction rises. This demonstrates how the CS coating’s compactness will be impacted by the amount of SP particles present. Additionally, if the SP ratio rises, the hardness also improves, as illustrated in Figure 13 [66].

Figure 12

Microstructural behaviour of the CSed and SP of Al-coated LA43M deposit [65]. (a and b) Surface morphology, (c and d) cross-sectional microstructure, and (e and f) etched cross-sectional microstructure of pure Al coated LA43M and shot-peened pure Al coated LA43M, respectively. (g and h) High-magnification SEM image of shot-peened Al coating and elemental distribution maps of coating/substrate interface.

Figure 13

Hardness characteristics [66].

Therefore, by following the above method, it is able to attain the desired mechanical, tribological, and electrochemical characteristics of cold sprayed coatings after SP. In summary, SP surface modification of the coating is done by compressive stresses. It reduces the surface roughness of the coating and will also harm the surface of the coating because of peening effect.

2.4 LRM treatment

The effect of LRM on cod sprayed coatings is demonstrated in Table 4. Astarita et al. [39] investigated the CS titanium coatings with post-LRM. They revealed that pores on the upper surface of the substrate were removed. The enhancement in the coating density at the LRM zone was owing to its melting and porosity moving to the free surface and following solidification. In the LRM, the surface treatment is developed with various zones: they are remelted, heat affected, and BM as seen in Figure 14. In the BM zone, fin lamellae were formed in the cold sprayed coatings, while coarser lamellae were developed in the heat-affected region (HAZ). They help to produce the growth of grains in the coating structure. The same phenomena were obtained by Rubino et al. [58]. The transition region is developed between the BM and the HAZ. The alpha grains were formed in the re-melted region because of sudden cooling. Likewise, the transition region was formed between the remelted and the heat-affected zone. Khun et al. [40] investigated the influence of laser power on the CS Ti64 deposit. They demonstrated that using high laser energy on the surface of the coatings causes many flaws, such as shrinkages and cracks, to form in the coatings due to sudden cooling. It is the sudden cooling after LRM that results in the enhanced hardness, wear, and corrosion resistance. The cracks produced in the coated surface may affect the deposition properties of various other coating materials after sudden cooling, as well as optimizing the LRM process parameters such as laser power, spot diameter, and traverse speed to attain the desired coating properties.

Figure 14

(a) Optical micro-graphical view of the cross-section of CS Ti deposit after LRM [16]. (b–f) SEM micro-graphical view of CS Ti deposit at various regions: (b) BM, (c) HAZ, (d) transition region between BM and HAZ, (e) remelted region (RZ), and (f) transition region between HAZ and RZ [39].

Kang et al. [56] examined the behaviour of aluminium and silicon composite coatings before and after LRM. In the cold sprayed condition, the bonding between Al and Si is lower because the mechanical interlocking does not take place between the particles. Due to the heat energy produced by LRM, the Si particles are melted completely and have a greater diffusion between the Al and Si particles. This results in enhanced properties of the coatings. Poza et al. [57] examined the hardness behaviour of CS sprayed Inconel 625 deposit and the impact of various laser scanning speeds. They revealed that the LRM depth was reduced as the scanning speed increased owing to fewer heat inlets. Furthermore, after LRM, the porosity of the deposit was considerably lowered, resulting in a greater coefficient of elasticity. By raising the laser scanning speed, the modulus of elasticity of the laser remelted deposit stayed unchanged. Moreover, the laser remelted deposit had a very low hardness value owing to the existence of columnar dendritic microstructural behaviour. Marrocco et al. [16] investigated the anticorrosion behaviour of the titanium coatings on the carbon steel. The open circuit and corrosion potential rates were shifted along the noble position after LRM, and the potentiodynamic polarization current (Ipp) decreased significantly. These findings showed that the laser remelted deposit showed greater anti-corrosion resistance compared to the pure titanium deposit. This is owing to the denser laser remelted layer and the defensive surface oxide layer. In summary, LRM may increase the density of the coatings. It results in a superior improvement in the tribological and electrochemical properties of the post-treated coatings, but there will be larger pores between the remelted layer and the substrate.

Sova et al. [67] identified that the LR operation significantly altered the mechanical behaviour of the CS deposit as illustrated in Figure 15. When compared to the CS deposit, the hardness of the laser-remelted region has greatly increased. As a result, the porosity of the deposit is massively diminished. After LR, the coated region is compressed by heat and also greater compressive residual stress raises the coated region’s hardness. In certain cases, the development of additional phases in the melting region also raises the hardness of the deposit. The tribological properties are also enhanced as the deposit hardness rises [55].

Figure 15

Microhardness mapping of the laser-treated deposit [67].