Plasma Sprayed Metal-Ceramic Coatings and ... - Springer Link

JTTEE5 20:927–938 DOI: 10.1007/s11666-011-9646-9 1059-9630/$19.00 ASM International

O.P. Solonenko, V.E. Ovcharenko, Yu.F. Ivanov, and A.A. Golovin (Submitted January 2, 2011; in revised form March 4, 2011) Composite powder obtained from mechanically crushed titanium carbide—metal binder cermet compacts deserves special mention for plasma spraying of wear-resistant coatings. However, cermet coatings sprayed using this powder have comparatively high porosity. The porosity causes the mechanical strength of the coating to largely deteriorate, and it also lowers the strength of the bond between the coating and the substrate. Computational and physical experiments were performed in this area to reveal the possibilities offered by pulsed electron beam irradiation for structural modification of 70 vol.%TiC-(Ni-Cr) powder coatings. The authors evaluated optimal values of process parameters for suitability in implementing a controlled thermal treatment of coatings under conditions of solid-liquid interaction of components in the cermet composition with each other and with the steel substrate. Evolution of the structure and physical properties of the cermet coatings under rapid heating and following cooling in a wide range of temperatures typical of pulsed irradiation conditions have been examined.

Keywords

finite-element modeling, metal-ceramic powder, plasma coating, post-treatment, pulsed electron beam, splat

1. Introduction Plasma spraying of coatings from powder materials, a process that can be implemented in a broad range of thermal and gas-dynamic parameters of plasma jets, offers a wide range of possibilities in application of protective and reinforcing coatings on surfaces (Ref 1-3). The use, in plasma spray process, of composite powders with particles formed by very hard inclusions in a metal alloy matrix presents a promising strategy in synthesis of wear-resistant coatings. Presently, composite powder materials are available that can be obtained by many methods, including chemical-vapor, galvanic, or gas-phase deposition of pure metals clad onto various particles (Ref 4); conglomeration of metal or nonmetal powders (Ref 5); mechanical alloying of particles (Ref 6); plasma synthesis of granulated metal-ceramic (cermet) powders formed by a thermoreacting powder mixture of substances (Ref 7); and hightemperature self-propagating synthesis (SHS) of highmelting compounds from a mixture of pure elements with a metal binder (Ref 8). O.P. Solonenko and A.A. Golovin, Khristianovich Institute of Theoretical and Applied Mechanics SB RAS, Novosibirsk, Russia; V.E. Ovcharenko, Institute of Strength Physics and Material Science SB RAS, Tomsk, Russia; and Yu.F. Ivanov, Institute of High-Current Electronics SB RAS, Tomsk, Russia. Contact e-mail: [email protected].

Journal of Thermal Spray Technology

Composite powders, obtained from mechanically crushed titanium carbide/metal binder cermet compacts synthesized by thermal explosion of starting elements under high-pressure conditions, deserve special mention. In cermet powders, particles have a finely dispersed inner structure formed by 1-3 lm equiaxed carbide inclusions uniformly distributed, throughout the whole particle volume, in a metal binder (Ref 9, 10). The high carbide phase inclusions volume content in the cermet powder particles, 50-80%, provides for a high viscosity of the binder melt with high-melting inclusions suspended in it. In addition, during SHS and plasma spray processes the carbide inclusions in individual cermet particles may undergo sintering, which results in the formation of a solid ultrafine-particle skeleton. All those processes define a low value of deformation of the particles impinging on the substrate or on previously deposited coating layers. That is why cermet coatings plasma sprayed from powder materials have comparatively high porosity at the interfaces between individual cermet splats and at the coating/substrate interface. The above porosity causes the mechanical strength of the coating to largely deteriorate, and it also lowers the strength of the bond between the coating and the substrate. Porosity in a plasma sprayed cermet coating can be reduced, with simultaneous strengthening of the coatingto-substrate bonding, by giving the coating a thermal treatment. The problem here consists of the necessity to heat the coating to a desired temperature without seriously affecting the metal substrate. In practice, such a high-temperature treatment of powder coatings can be organized using pulsed laser, ion beam, or electron beam irradiation (Ref 11-13). Of special interest here is highenergy treatment of cermet alloy with pulsed electron

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Plasma Sprayed Metal-Ceramic Coatings and Modification of Their Structure with Pulsed Electron Beam Irradiation

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beams, capable of heating a surface layer to high temperatures at depths of 0.1 mm (Ref 14). In the present paper, we report on computational experiments that were performed with the aim to reveal the possibilities offered by electron beam irradiation in modification of the structure of cermet powder coating TiC-(NiCr). To this end, we examined the evolution of the structure and physical properties of plasma sprayed cermet coatings under rapid heating and following cooling in a wide range of temperatures typical of pulsed irradiation conditions.

In the present physical model, the following assumptions were adopted. First, we assume that the energy flux is distributed uniformly over the beam cross section. According to the additivity principle and without regard for material porosity, the effective values of the density, the specific heat, and the heat conductivity in the cermet coating can be represented as q ¼ vTiC qTiC þ ð1 vTiC ÞqNiCr c ¼ vTiC cTiC þ ð1 vTiC ÞcNiCr k ¼ vTiC kTiC þ ð1 vTiC ÞkNiCr

2. Experimental The composite powders used for plasma spraying of cermet coatings were obtained by mechanical crushing of a SHS compact synthesized under high-pressure conditions. The composition of the compact was TiC-(Ni20 at.%Cr) at 70 vol.% of carbide inclusions. The cermet coatings were sprayed onto steel substrates using a plasma torch with interelectrode insert (Ref 15). Argon was used as the working and transporting gases. The coatings were sprayed from powders TiC + 30 vol.%Ni-Cr of 50-63 lm, 63-80 lm, 80-100 lm fractions. The operating conditions of the plasma torch were: working gas mass flow, 0.85 g/s; transporting gas mass flow, 0.06 g/s; arc current, 200 or 250 A, respectively, at arc voltages 102 and 111 V; and effective thermal power of the plasma jet at the nozzle exit plane of the gun, 9.4 or 12 kW. The spraying distance was 120 mm. For pulsed electron beam irradiation of the coatings, we used an electron beam assembly characterized by diameter of 1-3 cm and beam current of up to 300 A; the pulse duration of electron beam could be varied in the range 10-200 ls at an accelerating voltage of up to 25 kV, and the energy flux density in the beam of up to 100 J/cm2 (Ref 16, 17). The irradiation conditions under coatings treatment were: energy flux density in the electron beam, 40 J/cm2; pulse duration of electron beam, 150 ls; total number of electron beam pulses, 50, 100, 150, or 200; and pulse repetition frequency 0.3 Hz. The morphology and microstructure of the sprayed coatings were investigated using optical metallography and scanning electron microscopy. The elemental composition of the coatings was examined by the local x-ray spectrum analysis method.

3. Preliminary Numerical Analysis of the Electron Beam Treatment Process 3.1 Statement of the Boundary-Value Problem Numerical experiments were carried out using the software package described in Ref 18 and 19. The package was developed for modeling, by the finite-element method, of nonstationary boundary-value heat-transfer and phasetransition problems usually met in the field of treating surfaces and coatings with concentrated energy fluxes.

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Here, q, c, and k are the density, the specific heat, and the thermal conductivity, and vTiC = 0.7 is the volume fraction of TiC in the cermet coating. Second, we assume that at the melting point of the Ni-Cr alloy the TiC particles always remain solid unless the melting point of TiC is reached. The behavior of the TiC-(Ni-Cr) particles at higher temperatures is not considered in the present study since in the pulsed electron beam treatment of cermet particles it is normally required to provide for melting of the Ni-Cr component of the composition only, without reaching the melting point of titanium carbide. Then, the melting point of the coating can be defined as Tm = Tm,NiCr, and the melting heat, as Lm ¼ ð1 vTiC ÞLm;NiCr where Lm,NiCr is the melting heat of the Ni-Cr alloy. In calculations, as the substrate material for spraying the cermet coatings, steel SUS430 was chosen. The calculations were performed with due allowance for porosity P. The porous material was modeled by applying a proper correction to the effective values of q, c, k, and Lm according to ap = (1 P)a, where a is a property of dense composite, and ap is the same property of porous material. The thermophysical properties of involved materials and porous composites of different porosities are given in Table 1. Figure 1 shows a schematic of the calculation domain. The subscripts f and b refer to the coating and substrate materials; the subscript s, to the solid state of coating; and the subscript l, to the state of coating with melted Ni-Cr binder in the TiC-(Ni-Cr) composition. The heat transfer in the region X ¼ Xl [ Xs [ Xb is governed by the heat conduction equation qi ci

@T ¼ divðki grad T Þ; @t

i ¼ f ðs and lÞ; b

ðEq 1Þ

where qi, ci, and ki are the density, the heat capacity, and the heat conductivity of the ith material. In what follows, the material subscripts are used to refer to a particular material only in nonobvious situations. The boundary conditions were: Effect due to the external heat flux: @T k ¼ q @n Cq

ðEq 2Þ


Material TiC (Ni-Cr) 70%TiC + 30%(Ni-Cr)(a) (P = 0) 70%TiC + 30%(Ni-Cr)(a) (P = 0.3) 70%TiC + 30%(Ni-Cr)(a) (P = 0.5) SUS430

q, kg/m3

c, J/kg Æ K

k, W/m Æ K

Tm, K

Te, K

Lm, J/kg

4920 8340 5946 4162.2 2973 7700

572.5 460 538.75 377.13 269.38 585

17 12.2 15.56 10.9 7.78 26

3530 1663 1663 1663 1663 1767

4573 3653 ÆÆÆ ÆÆÆ ÆÆÆ ÆÆÆ

1.4 9 106 298,851 89,655.3 62,758.71 44,827.65 250,000

(a) Volume content of component in the cermet composition

q Ωl

Γsl

Ωs

Γ

Ωb

T

Γq

Tm

ξ*

Fig. 2 Correlation between the position of the melting front and the temperature

Γ0

Fig. 1 Calculation domain

Here, n is the unit vector of the outer normal to the boundary. Melting: @T @T dn þ k ¼ Lm q ; TjCsl ¼ T m ðEq 3Þ k @n Cslþ0 @n Csl0 dt Perfect thermal contact between the porous composite coating material and the substrate: @T @T k ¼k ðEq 4Þ @n Cfbþ0 @n Cfb0 Absence of heat transfer into ambient medium at the boundary C0: @T k ¼ 0 ðEq 5Þ @n C0 The initial conditions at t = 0 are Tðx; yÞ ¼ T0 ;

ðx; yÞ 2 X

ðEq 6Þ

The formulated boundary-value problem in Eq 1-6 was solved by the finite-element method using a time-implicit approximation. Note that, at the boundary Csl, boundary conditions of the first and second kind are set simultaneously. As a rule, first-kind boundary conditions are used for finding temperature fields, while second-kind boundary conditions are used for determining the motion dynamics of phasetransition fronts. The involvement of boundary conditions


ξ

of the second kind entails substantial difficulties related with the necessity to differentiate obtained solutions since the calculation of derivative ¶T/¶n in the vicinity of boundary defines the determination accuracy for front position. In the present study, we used an approach based on the fact that, in the finite-element method, first-kind boundary conditions in solving the heat-transfer problem can be rather naturally taken into account, while the second-kind boundary conditions can be used for determining the exact position of phase-transition fronts. For instance, in determination of the exact position of the melting front, the following reasoning can be used. With an overestimated value of melting-front velocity, there will be a greater amount of melted substance, and, hence, a greater amount of the heat spent for the phase transition is to be assumed; as a result, the temperature at the solid/liquid interface will assume a value below the melting point. On the contrary, with an underestimated value of melting-front velocity, the heat flux will remain uncompensated, causing melting-front overheating. Thus, the deviation of melting-front temperature from a given melting point of Ni-Cr inclusions in the cermet coating can be expected to vary in proportion to the difference between the melting-front coordinate and the exact value of this coordinate (see Fig. 2). In calculating the quantities of interest at the next time layer, the value of the melting-front coordinate can be refined using the iterative procedure: yðiþ1Þ ¼ yðiÞ f ðiÞ

yðiÞ yði1Þ f ðiÞ f ði1Þ

ðEq 7Þ

Here, the function f ðiÞ ¼ TðyðiÞ Þ Tm presents the local residual of the current melting-front temperatures in the iterative process. As zero approximation y(0),

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Table 1 Thermophysical properties of materials

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the authors used the front position at the previous time layer, while for the first approximation (i = 1) they used either an estimate based on the front velocity at the previous time step or a sufficiently small trial perturbation.

where (¶T/¶z)z=n+0 is the temperature gradient at the solidification front from the side of the melt at time t, and dn/dt is the solidification-front velocity. Then, with ts being the total solidification time of the melted layer in the coating, we can evaluate the cooling rate of the melt as

3.2 Numerical Data and Their Discussion Initially, a series of numerical experiments was performed for coatings of thicknesses D = 50, 100, and 1000 lm treated with heat fluxes in a broad range of q, from 5 9 107 to 1010 W/m2, which could be obtained with our electron beam facility (Ref 20). Calculations were carried out for conditions of heating of the coating surface over the melting point of the Ni-Cr binder in the composite particles; they were terminated if the melting point of TiC was reached at the coating surface. Figure 3 shows dependences that illustrate the effect the heat flux density has on the depth to which the (Ni-Cr) binder becomes melted by the moment the coating surface reaches the melting point of titanium carbide in cermet coatings of different thicknesses (50, 100, and 1000 lm). It is seen that in the examined range of the heat flux density for a 1000 lm thick coating this intensity turns out to be insufficient for complete melting of the (Ni-Cr) component throughout the whole coating thickness at the moment the surface temperature reaches the melting point of TiC. In a cermet coating 100 lm thick the (Ni-Cr) component can be melted throughout the whole coating thickness at the moment the surface temperature reaches the melting point of TiC in the range of heat flux densities from 5 9 107 to 108 W/m2. For a 50 lm thick cermet coating, the latter range is somewhat wider, from 5 9 107 to 5 9 108 W/m2. It should be noted that at heat flux densities in excess of 109 W/m2 the depth to which the (NiCr) binder becomes melted is completely independent of coating thickness, and it rapidly decreases with increasing the heat flux density. According to Ref 21, the current value of cooling rate during melt solidification can be represented as @T dn _ TðtÞ ¼ @z z¼nþ0 dt

1 T_ ¼ ts

Zts

1 _ TðtÞdt ¼ ts

0

Zts 0

@T @z

dn dt z¼nþ0 dt

ðEq 8Þ

Figure 4 shows the rate of melt cooling versus the heat flux density. As in the previous figure, it is seen that at high heat flux densities the coating thickness does not alter the general picture of the thermophysical processes, proceeding in this figure within a narrow near-surface layer that is much thinner than the coating. Let us illustrate the potential offered by the pulsed electron beam irradiation method in bringing the Ni-Cr component into melted state throughout the whole volume of the coating. We express the electron beam pulse energy in terms of the energy flux and pulse duration as Ei = q Æ ti. Then, with known energy-flux values and with the corresponding times required for heating the coating surface to the melting points of Ni-Cr and TiC components, in the coordinates Ei ti, we can determine the range of irradiation conditions under which, following one pulse, the coating surface would be heated to a temperature over the melting point of Ni-Cr yet below the melting point of TiC. Figure 5 shows the energy flux value (during electron beam pulse) that will ensure, after a single pulse, heating of the surface of 50 lm thick cermet coatings of porosities 0, 30, and 50% to the melting points of Ni-Cr and TiC. It can be concluded that, as the porosity of a plasma sprayed cermet coating increases, the heating depth of the coating under pulsed electron beam irradiation also grows in value. The heating depth also increases markedly with the increase in the duration of electron beam pulses (Fig. 6). As a result of the performed numerical experiments, it was established that the optimal range of regime parameters of the experimental setup (pulse energy and pulse duration) for treating the cermet coatings under 1E+12

T, K/s

1E+11 1E+10 50 μm

1E+09

100 μm

1E+08

1 mm

1E+07 1E+06 Fig. 3 Dependences illustrating the effect the heat flux density has on the depth to which the (Ni-Cr) binder becomes melted by the moment the coating surface reaches the melting point of titanium carbide in cermet coatings of different thicknesses (50, 100, and 1000 lm)

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1E+05 1E+07

q, W/m 2 1E+08

1E+09

1E+10

1E+11

Fig. 4 Cooling rate of the Ni-Cr-binder melt vs. heat flux density after the end of the electron beam pulse


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Fig. 5 Electron beam energy fluxes providing, following one electron beam pulse, for heating of the surface of 50 lm thick cermet coatings of different porosities (1, 10 —0; 2, 20 —30 vol.%; 3, 30 —50 vol.%) to the melting points of Ni-Cr and TiC

Fig. 7 Surface that defines the range of q (W/m2) in the electron beam within the operating space of regime parameters of the experimental setup: pulse duration ti, ls; pulse energy Ei, J/cm2. Different intervals of q are indicated with different colors: 1, [107;108]; 2, [108;109]; 3, [109;1010]; 4, [1010;1011]

prevailing influence on the formation of temperature profile in coatings under electron beam irradiation and, second, that one can exert control over the duration of the interphase interaction of cermet components under certain restrictions imposed on the range of temperature variation in the coating. Numerical estimates of optimal values of process parameters allow use of pulsed electron beam irradiation for modification of structure of cermet coatings under conditions of solid-liquid phase interaction of cermet components with each other and with the steel substrate.

Fig. 6 Melting depth of the Ni-Cr component in the 50 and 1000 lm thick cermet coatings vs. electron beam pulse duration (solid curves, dense coating; P = 0)

4. Experimental Results and Their Discussion

consideration involves the range of heat flux densities 107 < q < 5 9 108 W/m2. In Fig. 7 this range refers to the area on the surface q = q(ti, Ei) denoted with No. 1 and in part No. 2. Figure 8 shows the calculated curves of surface temperature and depth of heating of the coatings versus the number of electron beam pulses (the surface temperature here is assumed registered at the end of the pulse). As the number of electron beam pulses increases, the coating surface temperature grows in value; this temperature varies in proportion to coating thickness. The latter effect is defined by the influence of the substrate, whose thermal conductivity is much greater than that of the coating. That is why the depth of heating of the cermet coating to the melting point of the Ni-Cr component depends weakly on the number of electron beam pulses given to the sample: on increasing the latter number from 1 to 10 the melting depth increases by less than 1 lm. The data obtained in our computer experiments show, first, that it is the energy flux due to the electron beam and the duration of electron beam pulses that exert a

Figures 9(a) and (b) show the appearance of a large composite powder particle obtained from a mechanically crushed cermet composite synthesized under high-pressure conditions (Fig. 9a) and the microstructure of the particle (Fig. 9b). The particle consists of carbide inclusions of predominantly round shape with mean sizes 3-5 lm. After grinding, the obtained powder particles were sieved into three fractions, which were subsequently used for plasma spraying of cermet coatings (see Fig. 10). Figure 11 shows a typical SEM image of the surface of an as-deposited cermet coating. Along with high porosity, the coating features a heavily rough free surface. The coating, rather uniform in thickness in the as-deposited state, is formed by individual cermet-alloy splats with pores present at intersplat contacts. It is a well-known fact that, during plasma spraying, the coating is formed from splats successively laid onto the base (substrate or previously deposited layer) from individual powder particles. In this connection, of obvious interest is the morphology of individual splats deposited onto the polished surface of the steel substrate. Figure 12 illustrates typical splats formed from TiC + 30 vol.%Ni-Cr


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3200

T, K

3150 3100

100 μm 50 μm

3050

N 3000

1

3

5

7

9

(a)

(b)

Fig. 8 Temperature of heated coating surface (a) and the depth of heating of 50 and 100 lm thick cermet coatings to the melting point of the Ni-Cr component (b) vs. the total number of electron beam pulses given to the samples

Fig. 9 Appearance (a) and microstructure (b) of a composite powder particle: titanium carbide inclusions (dark areas) dispersed in a nickel-chromium binder (light areas)

Fig. 10 Size distribution of TiC-(Ni-Cr) cermet powder particles. Powder fractions are indicated that were subsequently used in plasma spraying of cermet coatings

composite particles, 45-50 lm in size, on their impingement onto smooth steel substrates. It is seen that, following impingement of such particles with a high volume content of solid titanium carbide inclusions finely dispersed in a liquid metal binder, a broad range of splat morphologies can be observed. A characteristic feature of the formation process of such splats is a low degree of deformation of the particles, resulting from a very high effective viscosity of the material ‘‘melt-solid dispersions.’’

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In heated composite particles, the binder melt presents no continuous liquid with solid inclusions suspended in it; instead, the melt is distributed throughout the particle volume in the form of submicrometer or micrometer inclusions. Such particles do not have enough time to undergo spheroidization in the plasma jet, and on impingement onto the substrate they normally disintegrate into separate fragments. The probability of particle disintegration, and also the size and number of formed fragments, depend on the effective size of individual melt inclusions inside the particle, and also on the mutual position of the inclusions. At a sufficiently large size of an individual liquid microvolume, the liquid may undergo disintegration due to its incompressibility and macroparticle deformation induced pressurization (see Fig. 12b, c, e, f) with flow out of the metal binder in the form of a film laid onto the substrate surface or in the form of individual, variously sized spherical droplets formed on disintegration of the melt film (see Fig. 12a, b, c, e, f). There forms a multitude of spherical, predominantly fine metal droplets (see Fig. 12a, c, d, f, g). Those droplets are formed by the forcing of flowing melt to the outer surface of the splat as the melt goes through the microchannels of different orientations formed inside the macroparticle on consolidation of titanium carbide particles in blocks carcasses.


Peer Reviewed Fig. 11 Surface morphology of a plasma sprayed cermet coating (SEM image). The particle sizes of the cermet powder from which the coating was sprayed were D = 63-80 lm

Fig. 12 Typical splats formed from cermet composite particles with a high density of hard dispersions and solid inclusions distributed throughout the volume of liquid metal binder

A preliminary qualitative analysis of the processes of splat formation from cermet particles impinging on the substrate shows that, along with the key physical


parameters such as the particle size, velocity, and temperature, the substrate temperature, and the state of substrate surface, the splat formation is substantially

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Peer Reviewed Fig. 13 Diagram of the experimental facility used for pulsed electron beam irradiation of materials in vacuum

influenced by the concentration of solid carbide inclusions, and also by the dispersion and mixing homogeneity of powder components in the material used in the synthesis of the cermet compact. Figure 13 shows the diagram of the experimental facility that was used for pulsed electron irradiation of materials. The facility comprises an electron gun, a control system, and a vacuum pumping system. While examining the effect of pulsed electron beam irradiation on the surface roughness and porosity of plasma sprayed coatings, we found that those characteristics in irradiated coatings, with all other conditions kept constant, were largely dependent on the dispersion of the initial composite powders. With variation of the dispersion, substantial modification of the surface of plasma sprayed cermet coatings was observed (see Fig. 14). Pulsed electron beam irradiation causes substantial modification of both the morphology and microstructure of plasma sprayed cermet coatings. In the course of irradiation, melting of the metal binder at the free surface and filling of pores with the melt proceed (Fig. 14a, b). With increase in the number of pulses, the melt fills the coating pores to a greater extent. On increasing the size of powder particles from 50 to 63 lm to 63 to 80 lm the intensity of heating and melting of a powder coating under pulsed electron beam irradiation grows in value. Following 200 electron beam pulses, almost complete melting of the metal binder in the surface layer of the coating is observed. Indeed, after 50 pulses the coating surface still exhibits a block structure resembling that in the initial coating, whereas after 200 pulses, the coating surface looks like a cast structure (Fig. 14c, d). This dependence of

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melting intensity of cermet coatings on the size of sprayed powder particles is even more distinctly observed in coatings sprayed from powder particles 80-100 lm in size: already after 50 pulses the coating surface becomes continuous, exhibiting a high characteristic luster (Fig. 14e, f). A closer examination of surface microstructure of irradiated cermet coatings showed that the surface layer in the coating was a polycrystalline aggregate with mean crystallite sizes largely defined by powder dispersion (Fig. 15). In the bulk of cermet coatings sprayed from 50 to 63 and 63 to 80 lm powders, as a result of electron beam irradiation there forms a structure with crystallite sizes varying within several tenth fractions of a micrometer (Fig. 15a, b). In a cermet coating sprayed from an 80 to 100 lm powder a cermet structure with nano- and submicrometer-sized composition components forms (Fig. 15c). The effect due to electron beam irradiation on the morphology and microstructure of coatings sprayed from composite powders of various dispersions can be traced by considering cross-sectional fractures of the coatings (see Fig. 16). With increasing the particle size of sprayed powder, the irradiated coatings undergo densification (the coating thickness decreases by a factor of 5-6), and at particle size of 80-100 lm a continuous (pore-free) coating forms (see Fig. 16c). The influence of pulsed electron beam irradiation on the physical properties of plasma sprayed cermet coatings is illustrated in Fig. 17 with the example of microhardness distribution across the coating in the direction from the coating surface to the steel substrate; here, data for samples irradiated with different numbers of electron beam pulses are given. The general behavior demonstrated by microhardness in the coating cross section is independent of the number of pulses: as we move farther from the coating surface, the microhardness decreases rather smoothly (by more than four times) to reach a minimum value at the steel substrate. On increasing the total number of pulses from 50 to 200, the microhardness at the coating surface decreases appreciably (by 20%). Yet, on moving from coating surface to steel substrate, the difference between the microhardness values decreases so that it vanishes at the coating/substrate interface. The reduced coating microhardness at the coating surface observed on increasing the number of irradiation pulses is a result of dissolution of the carbide phase in the metal binder, enhanced with increasing the total number of irradiation pulses given to the cermet coating. As we move from the coating surface to the interface with the steel substrate, the dilution of the cermet composition with iron melt becomes more and more a prevailing process, this process being independent of the size of sprayed powder particles. That is why right at the interface the microhardness values of the coatings sprayed from powders of different sizes turns out to be almost identical. The aforementioned characteristic features of the distribution of microhardness across the cermet coating show that a predominant influence on microhardness is exerted by the two processes: (i) on the coating surface, by the


Peer Reviewed Fig. 14 Surface morphology of cermet coatings plasma sprayed from composite powders of different dispersions after pulsed electron beam irradiation. Irradiation conditions: energy flux density, 40 J/cm2; duration of electron beam pulses, 150 ls; (a), (c), and (e) 50 pulses; (b), (d), and (f) 200 pulses; (a) and (b) D = 50-63 lm; (c) and (d), D = 63-80 lm; (e) and (f) D = 83-100 lm

Fig. 15 Surface microstructure of irradiated cermet coatings plasma sprayed from composite powders of different dispersions. (a) D = 50-63 lm, (b) D = 63-80 lm, and (c) D = 80-100 lm (energy flux density during an electron beam pulse, 40 J/cm2; duration of electron beam pulses, 150 ls; total number of pulses given to the sample, 200)


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Peer Reviewed Fig. 16 Microstructure of cross-sectional fractures of irradiated cermet coatings sprayed from composite powders of different dispersions. (a) D = 50-63 lm, (b) D = 63-80 lm, and (c) D = 80-100 lm (energy flux density during an electron beam pulse, 40 J/cm2; duration of electron beam pulses, 150 ls; total number of pulses given to the sample, 200)

carbide particles is insignificant, the mean size of the particles being 10-20 nm (Fig. 18).

5. Conclusions Our computer experiments showed that the energy

Fig. 17 Distribution of microhardness across a cermet coating plasma sprayed from 80 to 100 lm sized composite powder particles after irradiation of samples with different numbers of electron beam pulses. Irradiation conditions: energy flux density during an electron beam pulse, 40 J/cm2; duration of electron beam pulses, 150 ls. Number of pulses given to the samples: 1, 50; 2, 100; 3, 150; 4, 200 pulses

dissolution process of the carbide component in the metal binder melt and (ii) in the middle part of the coating and in the vicinity of the coating/substrate interface, by the dilution process of the cermet composition with the iron initially contained in the steel substrate. Indeed, a micro x-ray analysis of the elemental composition of the coating showed that after 200 electron beam pulses the Ni-Cr component undergoes remelting throughout the whole volume of the coating. It was found that on the coating surface there is a considerable amount of iron, which increases to 100% on passage from coating into substrate with concomitant reduction of the concentration of doping elements of the metal binder (Ni, Cr, Ti) to zero (see Fig. 18). Between the coating and the steel substrate, no sharp interface is observed; instead, the coating and the substrate form a diffusion zone of variable composition. In the surface layer of the coating, a high volume content of titanium carbide particles with sizes up to 0.5 lm was registered. In the middle part of the coating, the titanium carbide particles decrease in size, and the volume content of the particles decreases. In the immediate vicinity of the steel substrate, the concentration of

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flux density in the electron beam and the duration of electron beam pulses are parameters that have a predominant influence on the formation of the temperature profile in plasma sprayed coatings subjected to electron beam irradiation. Through variation of the total number of electron beam pulses and the time intervals between the pulses, one can regulate the duration of the interphase interaction of components in the cermet composition under restrictions imposed on the temperature profile in the surface layer of the coating. Optimal values of process parameters were evaluated, suitable for implementing a controlled thermal treatment of cermet coatings under conditions of solid-liquid interaction of components in the cermet composition with each other and with the steel substrate. Pulsed electron beam irradiation of plasma sprayed cermet coatings presents an efficient tool for rapid heating of powder coatings with their rapid quenching without affecting the metal base up to complete melting of metal binder in the coating. With increasing the particle size of plasma sprayed cermet powder and at irradiation conditions kept unchanged, the heating temperature of the coating increases. Pulsed electron beam irradiation of cermet coatings plasma sprayed onto steel substrates resulting in complete remelting of the metal binder leads to the formation, in the surface layer of the coating, of a high-dispersion structure formed by submicron titanium carbide particles dispersed in the metal binder based on the solid solution of nickel and chromium in iron. In the middle part of the coating, the aforementioned structure comes into a nanostructured state defined by the formation of a cermet structure from titanium carbide nanoparticles. In the immediate


Peer Reviewed Fig. 18 Distribution of the elemental composition of a cermet coating sprayed from 80 to 100 lm composite powder particles across the coating from free surface to steel substrate in electron beam-irradiated samples (energy flux density during an electron beam pulse, 40 J/cm2; duration of electron beam pulses, 150 ls; total number of pulses given to the sample, 200) and the microstructure of the cermet coating in its surface layer, middle part, and in the vicinity of the steel substrate (transmission electron diffraction microscopy data)

vicinity of the steel substrate the coating structure is defined by the formation of an iron-based solid solution with titanium carbide nanoinclusions.

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