Microhardness variation in heat-treated conventional and nanostructured NiCrC coatings prepared by HVAF spraying

Surface & Coatings Technology 203 (2009) 1406–1414

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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t

Microhardness variation in heat-treated conventional and nanostructured NiCrC coatings prepared by HVAF spraying Kai Tao a, Xianglin Zhou a,⁎, Hua Cui b, Jishan Zhang a a b

State Key laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China

a r t i c l e

i n f o

Article history: Received 18 June 2008 Accepted in revised form 19 November 2008 Available online 3 December 2008 PACS: 81.15.Rs 81.07.Bc 62.20.Qp 81.05.Bx Keywords: Thermal spray HVAF Nanocrystalline materials Nickel alloys Hardness Boiler tubes

a b s t r a c t A nanostructured NiCrC coating for high temperature erosion–corrosion protection was prepared by high velocity air-fuel (HVAF) spraying. While previous studies on nanostructured materials were most focused on the short-term development in microstructure and properties during heat treatment, no systematic investigation of the microhardness variation as a function of microstructural development during long-term heat treatment has been presented. In this work, HVAF-sprayed nanostructured NiCrC coating was heat treated at 650 °C for up to 200 h in air. The coating microstructures have been characterized by scanning electron microscopy, transmission electron microscopy, high resolution scanning electron microscopy and X-ray diffraction analysis. A Vickers microhardness tester was employed to determine the hardness variation of the coatings. In addition, conventional coarse-grained NiCrC coating produced by HVAF technique with the identical composition was also studied for comparison. The results indicated that the nanostructured NiCrC coating possessed a very compact and uniform microstructure, and exhibited good thermal stability during long-term heat treatment. The grain growth during thermal exposure caused the softening of the coating; however the carbide precipitation and content increasing resulted from phase transformation compensated the decrease by grain coarsening, and further led to the increase in the overall coating hardness. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Power station boiler tubes and other utility parts of coal-fired plants are subjected to frequent degradation by erosion–corrosion problems relevant to reliability and economics of these installations. The environment is characterized by high-temperature conditions together with aggressive atmospheres, leading to corrosive deposits adhere to the surface and to erosion processed due to the ash particles [1–3]. Attempts to decrease the maintenance costs of these components have increased interest in shielding them with protective coatings, such as thermal sprayed nickel- or iron-based alloy coatings [2–5]. High chromium content nickel-based alloys, as a widely used corrosion resistant coating material, possess good ability against oxidation, sulfur induced corrosion and hot corrosion [6,7]. A continuous Cr2O3 scale will form on the surface of the alloy with high bond strength when chromium content reaches the critical value, which protects the alloy from deterioration. However, the presence of erosive environments suggests that the application of harder, wear-

⁎ Corresponding author. Tel.: +86 10 82375385. E-mail address: [email protected] (X. Zhou). 0257-8972/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2008.11.020

resistant coating would perform better. Chromium has a strong ability to form carbide with carbon, which can effectively reinforce the matrix. Then Ni–Cr–C alloys with high corrosion and wear resistance together will be obtained only if the Cr/C ratio is optimized. Accordingly, a high Cr content Ni based alloy with addition of C element was selected in this study. Nanostructured (or nanocrystalline) materials, usually characterized by a microstructural length scale in the 1–200 nm range, have received considerable interest for the superior properties to their conventional coarse-grained counterparts, such as increased hardness and strength, improved ductility, enhanced corrosion and wear resistance [8,9]. Cryomilling, the mechanical attrition of powders within a cryogenic medium, is a recently developed method to synthesize nanostructured metallic or cermet powder [8,10]. The nanostructured bulk and coating materials using cryomilled powders as precursors have exhibited encouraging results in improving material properties [10,11]. The high velocity oxy/air-fuel (HVOF/ HVAF) process is one of the most popular thermal spraying technologies and has been widely adopted by many industries due to its flexibility, cost effectiveness and superior quality of the coatings produced [12]. The characteristics of ultra high flying velocity and relatively low temperature for injected feedstock powers much favor depositing of nanostructured coatings [9,13], as the nano-features of the materials could remain well after the spraying process. HVOF/

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HVAF-sprayed nanostructured coatings have proved great success in improving the performances of existing material systems, and are believed to be one of the most possible ways for commercial applications of nano-materials. Previous studies [14–17] on nanostructured materials were most focused on the microstructure and property evolution during short time (e.g. ≤8 h) heat treatment, aiming to provide references for preparation or transportation of these materials which were to serve at relatively low temperatures. However, industrially it is the longterm performance of NiCr–CrxCy coatings at elevated temperature that is significant. This has not been widely addressed, particularly for thermally sprayed nanostructured coatings. The purpose of this work was to study the long-term evolution in microstructure and hardness of the nanostructured coatings and hopefully to explore the potential application of nanostructured materials in high temperature environment. The response of HVAF-sprayed nanostructured NiCrC coating with cryomilled NiCrC powder as the feedstock was assessed after exposure at 650 °C for periods of up to 200 h in air. The conventional coarse-grained HVAF coating with the identical chemical composition using gas-atomized alloy powder as the feedstock was also investigated for comparison. 2. Experimental procedure Two types of NiCrC alloy powders, namely as-atomized powder with a conventional coarse-grained structure and cryomilled powder with a nanocrystalline structure, were employed as HVAF feedstock materials in this study. The former was prepared using gas atomization method with N2 as the atomizing gas. The latter was produced by ball milling a slurry of gas-atomized alloy powder in liquid nitrogen. The milling was carried out in a self-made attritor with a stainless steel vial and bearing steel balls (with 6.4 mm diameter) at a rate of 200 rpm. The ball-to-powder ratio was 25:1. After 20 h of milling, the cryomilled powder was removed from the attritor and used as feedstock powder for the following HVAF process. Specifications of both the as-atomized and cryomilled powders are listed in Table 1. An Intelli-Jet Activated Combustion High Velocity Air-Fuel (ACHVAF) spraying system of UniqueCoat technologies was used to deposit the conventional and nanostructured NiCrC coatings onto medium steel coupons under the same spray conditions, using propane as the fuel gas and nitrogen as powder carrier gas. Description of the HVAF spraying process has been reported [18]. Heat treatment was conducted at 650 °C in air. Samples were removed after 10, 30, 50, 100, 150 and 200 h. This temperature was used to provide the typical working temperature for most boiler tubes. Compositional analysis was performed by XRD in a RIGAKU DMAX-RB diffractometer using Cu-Kα monochromatized radiation. Metallographic sections were prepared from all coated specimens by standard metallographic procedures. Analysis on the cross sections was carried out using back scattered electron (BSE) scanning electron microscopy (SEM) (ZEISS SUPRA 55 equipped with EDS). Vickers microhardness measurements (LEICA VMHT 30 M Hardness tester) were conducted on the coating cross sections under a load of 300 g for 15 s. An average hardness was calculated from 20 indents per specimen.

Table 1 Nominal powder characteristics Powder type As-atomized Cryomilled

Cr

C

Chemical composition

Size

Ni

wt.%

d10/μm

d50/μm

d90/μm

52.26

45.88

1.86

80%NiCr–20%CrC

9.608 10.856

18.088 20.988

30.942 37.900

Note: d(number) denotes the particle diameter corresponding to the cumulative distribution of (number)%.

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Detailed microstructure analysis of the nanostructured coating was also performed by transmission electron microscopy (TEM) (HITACHI H-800) and high resolution electron microscopy (HREM) (JEOL-2010). TEM disc-shape specimen was prepared by mechanical polishing and ion-beam thinning of a flake-shape coating material, which was cut off from the substrate. The ion milling was carried out in a dual ion-beam miller at a current of 0.5 mA. 3. Results and discussion 3.1. Microstructural characterization of the as-sprayed coatings BSE images of the as-sprayed conventional and nanostructured NiCrC coatings are revealed in Figs. 1 (a)–(b) and 2 (a)–(b), respectively. The microstructure of the two coatings differed much from each other. For conventional NiCrC coating, some powders were heated to a relatively high temperature and fully softened during the spraying process; sometimes even a slight fusion occurred on the surface. This would lead to a layered morphology in the coating structure. Meanwhile, there were large quantities of powders which possessed relatively low temperatures and had not been adequately softened. Consequently these ones would retain much of their original shape upon deposition, often appearing sphere-like shape. According to the EDS analysis, chromium carbides exhibit dark-gray color in the images, and are labeled by arrow as region C. Oxides show black color in the image; however they are all distributed at the inter-splat boundaries and easily to be identified. It is noted that the carbides are heterogeneously distributed in the metallic matrix and the morphology varies in a large range, from coarse dendrite to fine round particles. Some of the splats are rich in carbide particulates, while others appear simply in dendritic morphology with bright color (metallic phase), or comprise from finely-eutectoid, lamellar structure exhibiting lightgray color in the BSE image. The feedstock employed for nanostructured NiCrC coating was cryomilled NiCrC alloy powder. During the cryomilling process, continuous welding and fracturing happened to the particles, which were caused by collisions between milling balls and the powder. Meanwhile, the microstructure also experienced significant changes, leading to the gradual refinement of grain size and finally the forming of nanocrystalline microstructure. Compared with the conventional counterpart, the HVAF sprayed nanostructured NiCrC coating has a more homogenous and denser microstructure. A lamellar characteristic could be seen from the BSE images; however, distinguished boundaries could hardly be established between the splats even under high magnification and this is beneficial to avoiding the possible intersplat bonding weakness in the coating. There are very fine carbide particles distributed in the metallic matrix, and it is difficult to differentiate individual tiny particulate under this magnification. A slightly heterogeneous distribution of carbide dispersion in the matrix generated the grayscale variation indicative of variation in the matrix composition, as shown in Fig. 2 (a). 3.2. Compositional evolution of the NiCrC coatings during heat treatment Compositional analysis by XRD (Fig. 3) indicated that all the carbides in the original as-atomized NiCrC powder and as-sprayed conventional NiCrC coating were in the form of Cr7C3. Based on the data in Cr–C binary phase diagram (from the software of Binary Alloy Phase Diagrams by ASM INTERNATIONAL), when the C content in this NiCrC alloy is larger than 4.5 wt.% with a melting temperature higher than 1576 ± 10 °C, should the Cr7C3 phase be likely to form. Considering the chemical analysis result of 1.86 wt.% for C content, which should result in the formation of Cr23C6, it is supposed that a heterogeneous distribution of C element happened in the alloy during the melting process before gas atomization. The so-calculated proportions of NiCr solid solution and Cr7C3 phase in the start NiCrC

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Fig. 1. BSE images of the microstructure of HVAF-sprayed conventional NiCrC coating: (a)–(b) as-sprayed state, after heat treatment in air at 650 °C for (c)–(d) 50 h and (e)–(f) 160 h.

alloy powder were about 80% and 20%, respectively, as listed in Table 1. This could also account for the heterogeneous microstructure of the as-sprayed conventional NiCrC coating. The oxide in the assprayed conventional NiCrC coating has a very low content, so the corresponding diffraction peaks were not clear in the XRD spectrum, where only NiCr (fcc) and Cr7C3 were identified. Whereas the XRD pattern (Fig. 4) of as-sprayed nanostructured NiCrC coating shows a complex composition, which contains NiCr matrix, Cr7C3, Cr2O3 and a metastable Cr2C phase. The metastable phase is believed to form in the rapid solidification of partly melted splat. The more extensive oxidation of nanostructured coating was thought to result from the cryomilling process and HVAF spraying process. It is evident from the XRD spectra of the heat-treated conventional and nanostructured coatings that microstructural transformations were well advanced during the thermal exposure. For the nanostructured NiCrC coating, the matrix in the as-sprayed state was a supersaturated FCC (face centered cubic structure) NiCr solid solution. As shown in Fig. 4, the XRD pattern included NiCr diffraction peaks with a distinct shift from the pure fcc Ni peaks towards lower diffraction angles, and relatively weak Cr7C3, Cr2C and Cr2O3 peaks. After treating for 10 h, the peaks for NiCr matrix became narrow and much sharpened, which was the result of grain growth

and lattice strain relaxation. And there was an obvious peak shift towards the higher diffraction angles when compared with the assprayed state. Oxidation occurred on the coating surface and led to the rising of Cr2O3 peaks. A new phase also appeared in the coating as Cr23C6, which is the most stable form of chromium carbide and is considered as the product of the interaction between Cr7C3 and solute Cr atoms in the coating. Besides, more characteristic peaks for Cr7C3 appeared in the profile as a result of exsolution, while the peaks for the Cr2C phase disappeared. As the thermal process continued, the NiCr peaks got increasingly sharper and shift slightly towards the higher angle direction, resulting from the crystal lattice parameters decrease during the exsolution process. However, the XRD spectrum varied in a very small extent after 10 h of heat treatment and it seemed that the exsolution process had approached an equilibrium state. Oxidation kept working on the coating surface during the treatment. Accordingly the Cr2O3 peaks grew stronger. The transformation from Cr7C3 to Cr23C6 also proceeded in the coating and gradually approached a stable condition, judging from the peak height ratio. It should be mentioned that the nanostructured coating sample after treating for 200 h was further processed, removing the tarnish layer that had formed on the coating surface by grinding, aiming to expose the underlying heat treated material that

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Fig. 2. BSE images of the microstructure of HVAF-sprayed nanostructured NiCrC coating: (a)–(b) as-sprayed state, after heat treatment in air at 650 °C for (c)–(d) 10 h, (e)–(f) 100 h and (g)–(h) 200 h.

was not influenced by oxidation. Consequently the oxide peaks were significantly reduced, both on the number and intensity, compared with those in the sample treated for 100 h. Nevertheless the phase composition was identical with that of the non-grinding coating specimen, with Cr23C6 as the major carbide form and the remained Cr7C3 phase as well. The composition evolution in the conventional NiCrC coating behaved in a manner similar to that of the nanostructured coating. The tarnish layer on the surface of the conventional coating heat

treated for 160 h was also removed by grinding. The treated conventional coating consisted of NiCr metallic phase, Cr23C6, Cr7C3 and Cr2O3. Like the case in the XRD spectrum of nanostructured coating treated for 200 h, the diffraction peaks of Cr2O3 were noisy and had low intensity, which indicated that this phase was present in a very low concentration. Besides the production of two new phases, the diffraction peaks seemed a little more defined and shifted towards the higher angle direction, when compared with those in the as-sprayed state, as shown in Fig. 3.

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Fig. 3. XRD spectra of the as-atomized NiCrC powder and conventional NiCrC coating in the as-sprayed state and heat treated at 650 °C for 160 h.

100 h of heat treatment, the numerous fine carbides had started to rearrange and bridge in each splat, revealing a tendency of being arrayed in the parallel directions. Till the time of 200 h of exposure, the trend presented in the coating earlier was more clear. Growth occurred primarily through strengthening of the inter-carbide bridges and continued carbide coalescence in the parallel orientations, forming a lamella-like structure with alternate layers of metallic matrix and stringer-like carbide that consisted of coalesced or bridged dispersions. The unique mechanism of carbide evolution in the nanostructured coating was thought to result from the cryomilling process of nanostructured NiCrC feedstock powder. During the milling process, steel balls and the under-milling powder continuously collided with each other. Consequently, the powder experienced severe deformation and underwent continuous fracture and welding, leading to the lamellar structure with large quantities of deformation bands in the cryomilled particles and thus the splats in the HVAF sprayed coatings. When the coating samples were exposed at elevated temperatures, these deformation bands in high energy state acted as the preferred region for carbide precipitation and other types of transformations.

3.3. Carbide precipitation and development during heat treatment

3.4. Microstructural development in the nanostructured NiCrC coating

For the conventional NiCrC coating, precipitation of fine carbides occurred throughout the coating during the initial 50 h of exposure (Fig. 1). As described in Section 3.1, the splats comprising the assprayed conventional coating could generally be sorted into three types in terms of microstructure — dendritic metallic phase, metalcarbide composite and lamellar eutectoid. The variation in the magnitude of solution content generated a variety of carbide distributions and morphologies with heat treatment. Among these kinds, metal-carbide composite splat had a high carbide density, and after heat treatment these splats were the preferential regions for precipitation; while the other two types of splat had relatively small amount of newly formed carbide particles. The dissolution mechanism of chromium carbide may occur by the C diffusion into the metallic matrix, because the diffusion coefficient of C in nickel (DC–Ni0 = 2.5 × 10− 4 m2 s− 1, 730–1020 °C) is higher than the Cr one (DCr–Ni0 = 6 × 10− 5 m2 s− 1, 850–1020 °C) [19]. Preferential nucleation occurred near those carbides from the feedstock, as the dissolved elements, mainly the C atoms, only diffused short distances from the parent carbides at this stage. However, the general coating presented a more uniform microstructure than the as-sprayed sample, with a large number of carbide distributed throughout the vision. Continued carbide precipitation and growth occurred out to 160 h, with the average particle size larger than that at the shorter exposure time. The fine carbide particle tended to coalesce and form larger carbides of complex morphologies. It was noted that some of the inter-splat boundaries exist as a preferred region for precipitation, leading to the surrounding of the splat by a carbide ring. Till then, the microstructure of the treated coating was much more homogeneous than as-sprayed coating, with carbide phase evenly distributed all over the coating area. All the splats presented the similar grey contrast, appearing entirely as a metal-carbide composite. The carbide evolution in heat-treated nanostructured NiCrC coating presented a much different mechanism. During the initial 10 h of thermal exposure, precipitation of very fine carbides occupied all regions of the coating. These particles dispersed uniformly in the matrix, as shown in Fig. 2. The uniform precipitation of carbide in nanostructured NiCrC coating was attributed to two aspects: first, the homogeneous distribution of Cr and C elements existing in the form of solute atoms or small carbide particles, which resulted from the cryomilling process; and second, the large quantities of grain boundaries in the nanostructured coating acting as quick diffusion channels, which promoted the diffusion process during the carbide precipitation. Beyond 10 h the carbides developed in a low rate. After

A plastically deformed metal or alloy contains large stored energy, and upon treating at elevated temperatures, it will normally revert to a lower energy state by structural evolution during recovery and recrystallization. Recovery is primarily due to changes in dislocation rearrangement and annihilation; while recrystallization may occur in which defect-free grains grow to consume the deformed or recovered microstructure [20]. Accordingly, the as-sprayed nanostructured NiCrC coating, as a non-equilibrium material, will undergo recovery, recrystallization, and subsequent grain growth during thermal exposure. TEM bright field images are displayed in Fig. 5, showing the representative microstructures of nanostructured NiCrC coatings during heat treatment. For the as-sprayed nanostructured coating, most areas showed a uniform nanostructure with most grains appearing equiaxed. The corresponding SAD pattern revealed polycrystalline rings, indicating that individual grains were separated by high-angle grain boundaries having random orientations with neighbors. The grain size measured from the TEM photos concentrated in the range of 30–50 nm, and the average value obtained by mean linear intercept method was 40.9 nm from 135 grains. The consistent distribution of grain size indicated that the nano-characteristics of both the metallic matrix and chromium

Fig. 4. XRD spectra of the nanostructured NiCrC coating in the as-sprayed state and heat treated at 650 °C for 10 h, 100 h and 200 h respectively.

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Fig. 5. TEM bright field images of the microstructure of nanostructured NiCrC coating in the (a) as-sprayed state with the inset showing the corresponding SAD pattern, and after (b) 10 h, (c) 100 h and (d) 200 h of heat treatment in air at 650 °C.

carbide from the cryomilled powder were well preserved through the HVAF process. During the exposure at 650 °C recrystallization and grain growth happened in the nanostructured NiCrC coating samples. After 10 h of heat treatment, recrystallization process seemed to be nearly completed. Plastically deformed grains with obvious faults (e.g. deformation twins) in the as-sprayed state had disappeared; meanwhile larger-sized faultless grains occupied most regions. The recrystallized grains had an equiaxed shape, with the grain size evenly distributed. After then, the microstructure of coating developed slowly with the continuing of treatment. The grains gradually increased to a bigger size while kept the equiaxed morphology. The whole process followed the general rule of normal grain growth. For each heat-treated coating sample, grain size was determined by measuring about 200 grains in the TEM photos using mean linear intercept method. The results of average values were plotted in Fig. 6 as a function of treating time. The data of grain size vs. treating time of the nanostructured NiCrC coating were further fitted mathematically, using a nonlinear, general

Fig. 6. Grain size development in HVAF-sprayed nanostructured NiCrC coating during heat treatment at 650 °C.

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curve fitting routine, to the popular equation of grain growth kinetics [21–23]: 1=n

d1=n −d0

= kt

where d is the mean grain size at time t, d0 is the initial grain size, k is the temperature-dependent rate constant, and n (≤0.5) is an empirical time constant determined experimentally. The best fitting curve was also drawn in Fig. 6 with the n value of 0.11. According to the curve, during the heat treatment at 650 °C grain growth occurred intensely at the initial 10 h, and then gradually stepped into a low rate evolution stage. After 50 h of exposure, the average grain size seemed to stabilize at about 100 nm. To further study the carbide transformation of nanostructured NiCrC coating during heat treatment, HREM was employed to characterize the coating specimen. Fig. 7 shows the HREM images of nanostructured NiCrC coating after 30 h of exposure at 650 °C. The typical microstructure consisted of equiaxed grains. Phase identification was conducted by FFT converting and calibrating of the atomic image of individual grains under high magnification. Due to the limit of resolution, this method was feasible only for carbide phases which had relatively large lattice parameters. Two forms of chromium carbide existed in the heat treated coating — Cr7C3 (marked with “A”) and Cr23C6 (marked with “B”), and were indicated in Fig. 7 (a). How-

ever, taking XRD results into consideration, the other non-marked grains are probably belonging to the fcc NiCr metallic phase. The similar morphology of different phases was well consistent with the TEM analysis. 3.5. Microhardness variation of NiCrC coatings as a function of heat treatment The similar general trend in microhardness variation occurred for both the conventional and nanostructured NiCrC coatings, as shown in Fig. 8. The hardness increased quickly in the initial period of heat treatment, and then gradually approached a stable value which was higher than the as-sprayed state. In all instances the nanostructured coating possessed a much higher hardness value. The hardness curve in the conventional coating varied in a simple manner, following the growing and stabilizing procedures, with the initial microhardness of 480.5 (HV300, 15 s) and the finally stable value of 546.0 (HV300, 15 s). When it comes to the nanostructured coating, the data developed in a more fluctuating way. The microhardness, with the initial microhardness of 697.1 (HV300, 15 s), increased rapidly over the initial 10 h, and then slowed the growing rate and gradually rose to the highest value after treating for 100 h. As the exposure continued, the hardness started to slightly decreased and finally approached a stable value of 728.1 (HV300, 15 s) after 150 h of exposure.

Fig. 7. HREM images of the nanostructured NiCrC coating heat treated at 650 °C for 30 h: (a) representative microstructure of the coating (“A” denotes Cr7C3 and “B” denotes Cr23C6); (b) detail of crystallite “A” in image (a) with its FFT pattern inset; (c) detail of crystallite “B1” in image (a) with the FFT pattern inset.

K. Tao et al. / Surface & Coatings Technology 203 (2009) 1406–1414

Fig. 8. Microhardness variation of the conventional and nanostructured NiCrC coatings heat treated at 650 °C as a function of time.

Grain size is known to have a significant effect on the mechanical behavior of materials, in particular, on the yield stress, σy, which, for materials with grain size d, is found to follow the Hall–Petch relation [24]: σ y = σ 0 + kd−1=2 where σ0 is the friction stress and k is a constant. If the workhardening is not large, the Vickers hardness, Hv divided by 3, usually approximates to the yield strength. Then the hardness also follows the Hall–Petch relation: Hv = H0 + kVd−1=2 this relation is applicable for most materials. The hardness usually decreases when the grains grow up during the high temperature exposure [15,25]. However it is not suitable for the hardness variation in the present study, where both the hardness and grain size increase after the heat treatment. Precipitation, depending on size, distribution and behavior of the precipitates formed, usually causes hardening. The as-sprayed NiCrC coatings had the supersaturated NiCr solid solution as its matrix. During the heat treatment, the dissolved Cr and C precipitated in the form of carbide, resulting in the decrease in the lattice constant of metallic matrix and thus the peak defining and shifting in the XRD spectra, seen in Figs. 3 and 4. High density of fine, hard and dispersed precipitates in the matrix, as shown in Figs. 1 and 2, would significantly increase the coating microhardness. Chromium carbide usually exists in three forms, namely Cr3C2, Cr7C3 and Cr23C6. Their relative thermodynamic stabilities are reflected by their standard free energies of formation. Of the three forms of chromium carbide, Cr23C6 is the most stable compound while Cr3C2 is the least stable. Owing to the fact that Cr23C6 is the most chromium-rich carbide, it is expected to be the equilibrium carbide phase in a low carbon and high chromium alloy. During high temperature exposure, carbide phases will undergo structural changes in the following manner: Cr3C2 → Cr7C3 → Cr23C6, accompanied by volume changes, which would affect the performances of material [26]. This also accounts for the carbide transformation from Cr7C3 to Cr23C6 in the present HVAF sprayed NiCrC coatings. According to the results of chemical analysis, the present NiCrC alloy has a C content of 1.86 wt.%. Assuming that all the C element has reacted with Cr to form Cr7C3, then the alloy will exist in the composition of 79.4%NiCr–20.6% Cr7C3, with the Cr content in the NiCr matrix of 34.14 wt.%. During the

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heat treatment, carbide transformation occurred significantly in the form of Cr7C3 → Cr23C6. If the entire C element exists in the form of Cr23C6, then the phase constituent of the present NiCrC material would be rewritten as 67.3%NiCr–32.7%Cr23C6. Therefore the volume content of hard phase in the NiCrC material will be much increased. This could also partly account for the increase of hardness in the HVAF-sprayed conventional and nanostructured NiCrC coatings after heat treatment. Because of the microstructural complexity of thermal spray coatings, a number of hardness mechanisms have been reported to be responsible for the changes in hardness with heat treatment. In the related studies of J. He et al. [14,27,28], both conventional coarse-grained and nanostructured 75%Cr3C2–25%NiCr coatings were synthesized by HVOF spraying. After isochronal heat-treatment in air for 8 h, microhardness of the conventional coating, with a initial value of 846 HV300, only increases slightly; while that of the nanostructured coating drastically increases from 1020 to 1240 HV300. The observed increase in microhardness is attributed to the precipitation of oxides, particularly in the nanostructured coatings, where the high density, nano-sized oxide particles lead to significant increases in the microhardness. Internal oxidation process was thought to be responsible for the precipitation of the dispersed Cr2O3 particles during heat treatment, which follows the chemical reaction: 2Cr3 C2 þ 17=2O2 ¼ 3Cr2 O3 þ 4CO2 : In the work of S. Matthews et al. [29], the microhardness response of HVAF and HVOF 75%Cr3C2–25%NiCr coatings was assessed following treatment at 900 °C for up to 60 days in air and argon. The air treated samples of both the HVAF and HVOF coatings were consistently harder than those treated in argon as a result of internal oxidation. And in the work of P.H. Suegama et al. [19], they also found an increase in the HVOF-sprayed 75%Cr3C2–25%NiCr coating after heat treatment in air at 760 °C and 880 °C, which resulted from the formation of chromium oxides and precipitation of chromium carbide. In the present study using NiCrC alloy powders, it is more likely that the matrix, rather than the carbide is preferentially oxidized according to the microstructure and relatively low carbide content (20 wt.%) of the feedstock powders. According to the results of XRD and thermal gravimetric data, no internal oxidation happened to the coating as a result of the excellent oxidation resistance of this alloy. To further verify the effect of internal oxidation on the coating hardness variation, the as-atomized and as-cryomilled NiCrC powders were heat treated in vacuum at 650 °C for up to 100 h. This was to exclude the possible oxidation caused by the thermal spray or heat treatment. Vickers microhardness measurement was undertaken for the heat treated powder samples using a load of 25 g for 15 s. A similar manner in powder microhardness variation was obtained for both powders. During the heat treatment, the hardness of both types of NiCrC powders increased initially with the exposure time, and then gradually approached a stable value which was higher than the hardness value in the as-sprayed state. The microstructural evolution was presented in terms of exsolution, recrystallization and grain growth in the metallic matrix, and the carbide precipitation, transformation, and its subsequent growing, same to that of the corresponding coatings. Therefore, internal oxidation mechanism is not accountable for the microhardness increase in the present HVAF sprayed NiCrC coatings. Also in the work of S. Matthews et al. [29], despite the significantly higher carbide content of the HVAF as-sprayed coating, its hardness is lower than the HVOF as-sprayed coating. This is in contrast to the results of Zimmermann and Kreye [30] who show a correlation between higher retained carbide content and higher hardness in Cr3C2–NiCr coatings sprayed by three HVOF techniques. S. Matthews et al. [29] attributed the difference in hardness of the as-sprayed coating primarily to the differences in matrix strengthening generated by carbide dissolution during spraying. Further, all samples showed a drop in hardness in the initial stage of thermal exposure due to the

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reduction in solid solution strengthening mechanism of matrix phase. And then hardness recovered and achieved stable hardness values after 30 days due to the precipitation and development of the carbide phase. Compared with HVOF spraying, one important characteristic of HVAF spraying is the relatively low temperature for the combustion gas, which allows the sprayed powder to deposit primarily in solid mode, thus effectively helps reducing the relevant problems, such as decarburization and oxidation [31]. Additionally, the carbide contents in the present NiCrC coatings are relatively low. Therefore, the dissolving of carbide phase in the NiCrC coatings during the HVAF spraying is negligible. The dissolved element in the supersaturated matrix of present NiCrC alloy is thought to originate from the gasatomization and cryomilling process. The carbide dissolution induced matrix hardening in the present case is less significant than that in the aforementioned, HVOF cermet coatings with high carbide content. This could account for the absence of hardness drop during the initial period of heat treatment. For the present HVAF-sprayed nanostructured NiCrC coating, the variation in hardness with heat treatment was the result of concurrent changes in microstructure — namely the precipitation and transformation of carbides and their subsequent growth, and the recrystallization and grain growth of the matrix. Recovery, recrystallization and grain growth of the metallic matrix occurred rapidly, reducing the deformation faults, increasing the grain size and relieving any internal stresses within the splats. These transformations led to the softening in the coating matrix. At the same time, nucleation of the carbide phases took place. The precipitation and transformation-induced content increasing in carbide phases not only compensated the softening by grain coarsening, but also increased the overall hardness of the coating during heat treatment. As for the conventional NiCrC coating, the hardness variation shares the similar mechanism with its nanostructured counterpart, but without the influence of notable grain growth, and is reflected by the simple evolving manner shown in Fig. 8. In addition, for thermally sprayed coatings, other factors, such as porosity, would also affect the overall hardness. Since both the present NiCrC coatings have very low porosities (b0.5%) which did not change evidently, it is not stressed here. 4. Conclusions In this work, the conventional and nanostructured NiCrC coatings were prepared by HVAF spraying, with gas-atomized and cryomilled alloy powders respectively as the feedstock. The microhardness variation of the coatings was studied as a function of microstructural development following heat treatment at 650 °C for periods of up to 200 h in air. Both the as-sprayed NiCrC coatings possessed a compact microstructure, especially for the nanostructured coating, which exhibited a more homogeneous morphology with uniformly distributed fine carbide dispersions and a much higher microhardness. During the heat treatment, recrystallization and grain growth occurred in the nanostructured coating, accompanied by exsolution, carbide precipitation, transformation and its subsequent development. The nanos-

tructured coating exhibited excellent thermal stability, whose average grain size stabilized at about 100 nm after 50 h of exposure at 650 °C. The microhardness variation in the nanostructured NiCrC coating was the result of concurrent grain growth and carbide development. Grain growth caused the softening of the coating; however, the carbide precipitation and content increasing resulted from phase transformation (Cr7C3 → Cr23C6) compensated the hardness decrease by grain coarsening, and further led to the increase in the overall coating hardness. The microhardness variation in the conventional NiCrC coating showed the similar mechanism to its nanostructured counterpart, but without the influence of notable grain growth. The results from this work demonstrated the nanostructured NiCrC coating as a hopeful candidate for boiler tubes protection during high temperature service. Acknowledgement This research was supported by the National High Technology Research and Development Program of China (“863” Program) (No. 2002AA331080). References [1] J. Stringer, Surf. Coat. Technol. 108–109 (1998) 1. [2] I. Fagoaga, J.L. Viviente, P. Gavin, J.M. Bronte, J. Garcia, J.A. Tagle, Thin Solid Films 317 (1998) 259. [3] T.S. Sidhu, S. Prakash, R.D. Agrawal, Thin Solid Films 515 (2006) 95. [4] T.S. Sidhu, S. Prakash, R.D. Agrawal, Scr. Mater. 55 (2006) 179. [5] D.J. Branagan, M. Breitsameter, B.E. Meacham, V. Belashchenko, J. Therm. Spray Technol. 14 (2005) 196. [6] S. Zhao, X. Xie, G.D. Smith, S.J. Patel, Mater. Chem. Phys. 90 (2005) 275. [7] B. Wang, J. Gong, A.Y. Wang, C. Sun, R.F. Huang, L.S. Wen, Surf. Coat. Technol. 149 (2002) 70. [8] J. He, J.M. Schoenung, Mater. Sci. Eng., A 336 (2002) 274. [9] T. Grosdidier, A. Tidu, H.L. Liao, Scr. Mater. 44 (2001) 387. [10] D.B. Witkin, E.J. Lavernia, Prog. Mater. Sci. 51 (2006) 1. [11] J. He, J. Schoenung, Metall. Mater. Trans., A 34 (2003) 673. [12] T.S. Sidhu, R.D. Agrawal, S. Prakash, Surf. Coat. Technol. 198 (2005) 441. [13] J. He, M. Ice, E. Lavernia, S. Dallek, Metall. Mater. Trans., A 31 (2000) 541. [14] J. He, E.J. Lavernia, Mater. Sci. Eng., A 301 (2001) 69. [15] M. Lewandowska, K.J. Kurzydlowski, Mater. Charact. 55 (2005) 395. [16] K.W. Liu, F. Mucklich, Acta Mater. 49 (2001) 395. [17] K. Chung, R. Rodriguez, E. Lavernia, J. Lee, Metall. Mater. Trans., A 33 (2002) 125. [18] K. Tao, J. zhang, H. Cui, X. Zhou, J. Zhang, Fabrication of conventional and nanostructured NiCrC coatings via HVAF technique, Trans. Nonferr. Met. Soc. China 18 (2008) 262. [19] P.H. Suegama, N. Espallargas, J.M. Guilemany, J. Fernandez, A.V. Benedetti, J. Electrochem. Soc. 153 (2006) B434. [20] F. Zhou, X.Z. Liao, Y.T. Zhu, S. Dallek, E.J. Lavernia, Acta Mater. 51 (2003) 2777. [21] T.R. Malow, C.C. Koch, Acta Mater. 45 (1997) 2177. [22] F. Liu, G. Yang, H. Wang, Z. Chen, Y. Zhou, Thermochim. Acta 443 (2006) 212. [23] J. Lee, F. Zhou, K. Chung, E. Lavernia, N. Kim, Metall. Mater. Trans., A 32 (2001) 3109. [24] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427. [25] W. Hu, M. Li, M. Fukumoto, Mater. Sci. Eng., A 478 (2008) 1. [26] G.Y. Lai, Thin Solid Films 64 (1979) 271. [27] J. He, M. Ice, E. Lavernia, Metall. Mater. Trans., A 31 (2000) 555. [28] J. He, M. Ice, E.J. Lavernia, Nanostruct. Mater. 10 (1998) 1271. [29] S. Matthews, M. Hyland, B. James, Acta Mater. 51 (2003) 4267. [30] S. Zimmermann, H. Kreye, in: CC Berndt (Ed.), Thermal Spray: Practical Solutions for Engineering Problems, ASM International, Materials Park, OH, USA, 1996, p. 147. [31] A. Verstak, V. Baranovski, Thermal Solution: Advances in Technology and Application, DVS — German Welding Society, Osaka, Japan, 2004, p. 551.