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Article

Development and Research of New Hybrid Composites with Increased Requirements for Heat and Wear Resistance

by
Peter Rusinov
1,*,
Chao Zhang
2,
Polina Sereda
1,
Anastasia Rusinova
1,
George Kurapov
1 and
Maxim Semadeni
1
1
Department of Engineering of Control Systems, Materials and Technologies in Mechanical Engineering, Kuban State Technological University, 350072 Krasnodar, Russia
2
School of Mechanical Engineering, Yangzhou University, Yangzhou 225127, China
*
Author to whom correspondence should be addressed.
Ceramics 2025, 8(1), 8; https://doi.org/10.3390/ceramics8010008
Submission received: 1 December 2024 / Revised: 29 December 2024 / Accepted: 16 January 2025 / Published: 18 January 2025
Browse Figures
Versions Notes

Abstract

Hybrid layered reinforced materials are able to increase the reliability, durability, and expand the functionality of high-temperature components in supercritical and ultra-supercritical power plants and in oil, gas, and petrochemical equipment operating under conditions with multifactorial influences (temperature, force, deformation). As a result of this research, surface reinforced ceramic composite materials with a gradient distribution of properties have been developed. These materials include thermal barrier layers (Gd2O3-Yb2O3-Y2O3-ZrO2) and Ni-based layers reinforced with ceramic carbide and oxide particles. They are strong, have a high heat and wear resistance, and provide the specified functional and mechanical properties. The formation technology for the hybrid composites has also been developed. This technology includes the mechanical alloying of powder compositions, which is followed by vacuum plasma spraying. The structure of the powder compositions and composite layers, the density of the obtained composite materials, and the heat and wear resistance of the composites have also been investigated. The microhardness of the alloy layers of the hybrid composite materials Hastelloy X–GYYZO–material 1 and Hastelloy X–GYYZO–material 2 was as follows: super alloy Hastelloy X, HV0.2 = 3.8–3.95 GPa; layer GYYZO, HV0.3 = 16.1–16.7 GPa; layer material 1, HV0.3 =18.3–18.8 GPa; layer material 2, HV0.3 =19.1–19.6 GPa. The influence of the refractory phase of HfC and TaC on the strength of the composites was studied. It was found that the maximum strength (710–715 MPa) in the composites Hastelloy X—GYYZO—material 1 and Hastelloy X–GYYZO–material 2 is achieved with a content of HfC and TaC–27–28%.

Graphical Abstract

1. Introduction

Functional Gradient Materials (FGMs) are known for their high functional and mechanical properties. These layered materials have different physico-chemical properties. As a rule, a change in the properties of FGMs is associated with a corresponding change in the chemical composition or physical structure of the material. These materials are mainly used in the production of components and devices operating under extreme conditions (large gradients of mechanical stresses and temperatures) [1,2]. FGMs increase the reliability and durability of modern machine components and mechanisms operating under conditions of high temperature and frictional–cyclic loading. Today, it is the only economically feasible way. The above applies, first of all, to the high-temperature components and assemblies of supercritical and ultra-supercritical power plants and the components and assemblies of oil, gas, and petrochemical equipment. Multicomponent, composite, and ceramic surface-modified layers are most widely used to increase the durability of units under high-temperature power exposure [3,4,5,6,7,8]. Such materials increase the durability of the units under high loads by 4–10 times. Much research is devoted to heat-resistant thermal protective layers, which can reduce the operating temperature of the product surface [9,10,11,12]. This opens up the possibility of further increasing the temperature and helps to increase the power and efficiency of the unit [13,14,15,16].
The method of vacuum plasma spraying (VPS) of metals and ceramic materials opens up certain prospects for new heat-resistant thermal protective surface-modified layers [17,18,19,20].
Applying surface-modified layers means not just increasing or improving the performance characteristics of a detail but creating a fundamentally new composition. This composition does not just have the characteristics of the base materials and protective layers; it obtains different properties that change the design of a detail or assembly and significantly increases their productivity and reliability.
New thermal barrier layers increase the operating temperature, which increases the power of the devices by 6–17%, depending on the operating conditions. This determines the task of developing new protection systems for units operating at high temperatures and voltages [21,22,23,24,25,26].
Based on this, there are important tasks to improve the existing methods and create new ones for surface-modified layers, to study their structure and properties, and to develop new material compositions.
To predict the behavior of new layered composites under extreme operating conditions, it is necessary to know their physical and mechanical properties (density, melting point, coefficient of thermal expansion, hardness). Therefore, it is advisable to obtain such composites in the form of thick layers, study their mechanical and physical characteristics, and then use the best of them in the form of protective surface materials.
The main requirement for heat-resistant materials is the ability to withstand the effects of oxygen at high temperatures. The physical phenomena that influence the heat resistance of composites include polymorphic transformations and recrystallization. Even surface-modified layers with zero porosity can lose their protective properties as a result of recrystallization. This promotes the penetration of gases through the layers to the base metal due to boundary diffusion [27,28,29]. An important means of increasing the heat resistance is to ensure their constant chemical composition. The chemical composition of the material can change, either as a result of the interaction with the gaseous environment or due to the interaction with the base material. The chemical destruction of the material, when interacting with gases, is prevented by the continuous gas-impermeable film in the coating–gas boundary layer. Such a layer is formed, for example, when molybdenum disilicide is heated in the air [30,31]. The same happens with NiCrAl alloys with an Al content of over 5% [32,33,34] when a glassy film of silica and α-Al2O3 is formed on the surface, isolating them from the gaseous environment. Sometimes, to prevent the migration of gaseous oxidizing atoms onto the surface of the protective coating, it is necessary to deposit a thin layer of glassy material with a high viscosity [35,36,37]. It is almost impossible to prevent the interaction of the protected material with the surface layer at high temperatures. In the case of such an interaction, there is mutual diffusion of the layer and base components, and then the protective and strength properties of the layer may be lost before the permissible period. Hence, the chemical composition of the base material and the formation technology are very important for the surface layers. As shown in [38,39], the reliability and durability of the surface composites largely depend on the chemical composition of the protected material. This was also established during the development and study of the diffusion properties for layers on niobium, tantalum, and their alloys [40,41,42,43]. Thus, the need for the long-term preservation of protective layers requires the absence of an interaction between the surface layer and the protected material. On the other hand, when forming protective layers of hybrid composites, to ensure the adhesion of the underlying layer to the protected surface, it is required to obtain a high diffusion rate. To combine these requirements, it is necessary to ensure that their interaction occurs only at the moment of fixing the surface layers. However, during the subsequent operation of the material, the speed of this interaction must be sharply reduced; otherwise, as already indicated, such composites quickly collapse. This can be achieved by introducing additives into the composite layers that reduce the diffusion processes or by creating barrier layers at the interface between the surface layer and the protected material [44,45,46,47].
When choosing a strengthening phase, one should proceed from its thermodynamic properties, its diffusion mobility in the matrix, and other parameters that take into account the operating conditions of the composition [48,49,50,51]. The volumetric content of the strengthening phase is determined by the technological capabilities to evenly distribute the components included into the alloy [52,53,54].
The problem of creating composite materials can be divided into two main stages which are to a certain extent interrelated. Firstly, the task is to develop the surface layers made of composites, which, in combination with the base material, should produce compatible and practically useful systems. Secondly, it is necessary to develop a technology for producing layered composites that would ensure their maximum reliability and durability during operation. The development of layered surface composites with a gradient of properties is extremely relevant for various industries (energy, oil and gas, mechanical engineering, etc.). The underlying functional layer (barrier layer), in contact with the base (protected part), should reduce the diffusion processes at the interface between the surface layer and the protected material. The overlying layer, in contact with an aggressive environment, should have increased heat resistance, wear resistance, and strength [4,39]. A correctly chosen sequence of layers provides high mechanical and functional properties for the composite materials operating at high temperatures [44,45].
The purpose of this work is to create surface reinforced layered composite materials with a gradient distribution of properties, including thermal barrier layers and layers of material with carbide and oxide inclusions, with high heat resistance, strength, and wear resistance, providing the specified functional and mechanical properties.

Research Objectives

To achieve the stated objective of the study, the following tasks were solved:
-
to justify the choice of developing new heat- and wear-resistant composite materials for the energy and oil and gas industries;
-
to justify the choice of equipment and technology for the formation of hybrid composite materials;
-
to assess the influence of the high-energy mechanical treatment (HMT) (mechanical alloying) of the powder compositions on the quality and mechanical properties of composites;
-
to study the structure of the powder compositions after the HMT, and to describe the mechanisms of interaction between the powder materials and the grinding media;
-
to study the structure and structural parameters of the composite materials;
-
to study and assess the density of the resulting composite materials (the effect of the content of the refractory ceramic phase on the porosity and average pore size of the composite layers);
-
to research and assess the mechanical properties of the composites (influence of the carbide phase content (HfC, TaC) in the composite layers on the flexural strength);
-
to study the heat resistance of the developed composites (material 1 (Ni-HfC-Re-ZrO2-Y2O3-Gd2O3-Yb2O3-Al2O3), material 2 (HfC-Re-TaC-ZrO2-Y2O3-Gd2O3-Yb2O3-Al2O3)); to compare the heat resistance of the developed composites with previously known ones (Ni-Al2O3 (24% wt.); NiCrY-Al2O3 (10% wt.); NiCrAl-ZrO2 (37% wt.)); and to assess the influence of the carbide and oxide phases on the heat resistance of the composites;
-
to study the wear resistance of the developed composites and to compare the wear resistance of the developed composites with previously known heat-resistant composites (NiCrY-Al2O3 (10% wt.); NiCrAl-ZrO2 (37% wt.)).

2. Materials and Research Methods

2.1. Materials

We chose the alloys Hastelloy X (NiCrFeMo) and Rene N6 (Haynes International, Kokomo, IN, USA), which are used in energy, oil and gas, and other industries, as the base [55,56,57,58]. We chose the ceramic material Gd2O3-Yb2O3-Y2O3-ZrO2 (GYYZO) as the thermal barrier layer adjacent to the base (Hastelloy X, Rene N6 alloy) [59]. The thermal barrier layer of ceramic material provides the necessary adhesion strength (chemical interaction) with the base material (Hastelloy X, Rene N6 alloy). During the subsequent high-temperature operation of the material, the diffusion interaction rate sharply decreases, which ensures the specified protective functions of the material. The combined ceramics material 1 and material 2 were developed as the surface heat- and wear-resistant layers (Table 1). A plasma jet was used as the source of the thermal energy. All this allows the formation of a fine-grained structure due to the short interaction time with the surface.
The content of the main chemical elements in the materials is presented in Table 1.

2.2. High Energy Machining

Before VPS in a protective atmosphere, high-energy mechanical treatment (HMT) was carried out. This process, in fact, was a mechanical alloying of the powder compositions. HMT improves the functional and mechanical properties of the surface. It was carried out in a water-cooled planetary centrifugal ball mill GEFEST-2 (AGO-2U) (Gefest, Moscow, Russia) with the following parameters: drum volume—150 cm3, ball-to-load ratio—(10–15):1, ball diameter—6 mm, ball material—WC, carrier rotation speed—890 rpm, drum rotation speed—1000 rpm. HMT was carried out in an inert argon environment. VPS in a protective atmosphere of the layered hybrid composites Hastelloy X/Rene N6 + GYYZO + material 1 and Hastelloy X/Rene N6 + GYYZO + material 2 was carried out after high-energy mechanical processing (mechanical alloying).

2.3. Formation of Layered Composites

Thermal barrier GYYZO and functional layers of material 1/material 2 composites were applied using the VPS method. Before VPS, the powder compositions were dried in a vacuum oven for 3–6 h at a temperature of 120–180 °C. To increase the adhesion strength of powder materials, the process included preliminary sandblasting of the samples’ surface, made of Hastelloy X and Rene N6 alloys.
The microstructure was studied with a super-high-resolution scanning electron microscope (JSM-7500F) (Akishima, Tokyo, Japan) and a transmission electron microscope (JEM-2100) (Akishima, Tokyo, Japan). Chemical analysis was carried out with an Arcmet 8000 (Tokyo, Japan) optical emission analyzer and on an energy-dispersive INCA x-sight attachment of a JSM-7500F electron microscope. The thickness was measured using a DMS 2E (GmbH, Kassel, Germany) and digital micrometer MKC-25 (Moscow, Russia) thickness gauge (sample preparation included grinding and polishing). X-ray phase analysis was carried out on a Dron-7M (Bourevestnik, Saint Petersburg, Russia) device in Cu-Kα radiation.
During the technology development for surface modification of Hastelloy X and Rene N6 alloys with materials (Table 1) using the VPS method, defects of cracking and peeling were not observed. This was achieved through the correct choice of chemical layered compositions, the calculation of linear thermal expansion coefficients, and modification technologies.
Since the formation of composite layers with the VPS method is multifactorial, algorithmic experimental planning was used for the technology. The main functional and mechanical parameters of the layers were as follows: adhesion, porosity, structure, and phase composition. The preliminary analysis demonstrated that the main VPS parameters affecting the structure and quality of the layers should include: arc current (I = 380 A), voltage (U = 40 V), flow and composition of plasma-forming gas, flow of powder and transport gas, spraying distance (150 mm), spraying angle (75°), moving speed and feed of the plasmatron, and rotation speed of the coated detail. The multifunctional technological complex (Figure 1a) allowed for ionic cleaning of the product before the formation of layered composites.

2.4. Properties

The microhardness was measured using a Falcon 503 (Brave Electronics Co., Limited, Hong Kong, China) using a diamond pyramid with an apex angle of 136° and an indenter load of 200–400 g. The porosity was determined by the method of hydrostatic weighing. Distilled water was used as the working fluid. Weighing was carried out on an “ADV 200” (Generator MSK, Moscow, Russia) analytical balance. To determine the density, we used plates of composite layers measuring 15 mm × 15 mm × 3 mm. After mechanical polishing and before weighing, the samples were degreased in carbon tetrachloride. The accuracy of measuring the density of composite layers was ±0.013 g/cm3.
Experimental density (ρE) was calculated using the formula
ρE = P1·(ρ1 + ρ2)/(P1 − P2) + ρ2
where P1, P2—weight of samples in air and water; ρ1, ρ2—density of water and air at test temperature.
The calculated density (ρC) of the material was calculated using the formula:
ρC = ρi1 · Cvi1 + ρi2 · Cvi2 + …+ ρii · Cvii
where ρi1—density of the first phase included in the material (g/cm3); ρi2—density of the second phase included in the material (g/cm3); ρii—density of i-phase included in the material (g/cm3); Cvi1—volumetric content of the first phase included in the material (%); Cvi2—volumetric content of the second phase included in the material (%); Cvii—volumetric content of phase i included in the material (%).
The relative density (ρR) of the material was calculated using the formula:
ρR = ρE/ρC.
Wear tests of materials were carried out on equipment 2070 SMT-1(Za-padpribor, Moscow, Russia).
Bend testing was carried out with Instron 8801 (lnstron, Norwood, MA, USA) equipment. These tests were carried out according to a three-point loading scheme, recording a “force-deflection” diagram at a loading speed of 1 mm/min. For this purpose, flat samples measuring 30 mm × 150 mm × 3 mm were used.
To test materials for heat resistance in calm air, a high-temperature SNOL (SNOL-TERM, Moscow, Russia) oven was used. Material samples measuring 10 × 15 mm2 were studied at a temperature of 1250 °C for 50 h. The kinetics of corrosion destruction was not studied; only the total weight loss and corrosion rate of the material were determined.
The coefficient of linear thermal expansion was measured using SDTA840 (Mettler Toledo AG, Moscow, Russia) equipment.
The coefficient of linear thermal expansion (α) was calculated using the formula
α = (LT − L0)/L0·(T − T0)
where LT—sample length at temperature T; L0—sample length at room temperature T0.

3. Results and Discussion

Mechanical alloying of powder compositions occurs in the field of three inertial forces: the Coriolis force and two centrifugal ones. The centrifugal forces acting on the balls and powder are many times greater than the force of gravity, resulting in the rapid activation of the powder. During the HMT process, the destruction of the oxide and ceramic particles occurs by micro-chipping; the powder particles acquire a rounded volumetric shape with a pronounced effect of mechanosynthesis. In the process of high-energy mechanical processing, the powder particles accumulate energy released during the VPS process. This increases the adhesive properties of the composite layers and the quality of the layers (i.e., reduces porosity).
The use of an inert medium reduced the content of oxides in the composite layers. This may have a negative impact on the functional and mechanical properties of the hybrid composite materials. The use of a multifunctional technological complex (Figure 1a) allows for ionic cleaning of the product before the formation of the layered composites. This helps to increase the adhesion strength of the layers to the base. The increase in the adhesive strength of the composite layers is explained by the combined effect of the sprayed material. It consists of VPS in an argon environment and the plastic deformation of the sprayed particles at the spraying temperature, accompanied by a high particle speed of the sprayed mechanically activated powder material, has a dynamic effect. The adhesive strength of the layered composite materials was as follows: Hastelloy X/GYYZO—32 MPa; GYYZO/material 1—34 MPa; Rene N6/GYYZO—29 MPa; GYYZO/material 2—31 MPa. The adhesive strength of the materials after HMT was Hastelloy X/GYYZO—39 MPa; GYYZO/material 1—43 MPa; Rene N6/GYYZO—34 MPa; GYYZO/material 2—38 MPa.
The compositions “material 1, material 2” (Table 1) were selected as materials of high heat resistance and wear resistance for the functional layers of the hybrid composite materials. These materials, unlike other superhard materials, are not subject to destruction at high temperatures and do not interact with the base Hastelloy X/Rene N6. To prevent the interaction of the chemical compositions of the materials “material 1, material 2” (Table 1) with the base, thermal barrier layers based on the GYYZO material were previously deposited. The chemical compositions of the materials were selected in such a way that the coefficients of linear thermal expansion of the materials at high temperatures were approximately the same. To increase the wear resistance and heat resistance, we added refractory carbide particles Ta and Hf to the functional surface layers based on Ni (Table 1). Nickel was chosen as the binding material.

3.1. Study of the Powder Composition and Structure After HMT

The advantage of HMT is the possibility of obtaining materials for coating with the required chemical and granulometric composition. The simplest way to obtain them is to mix the components. As a result of the “cold” intense plastic deformation that occurs during HMT, significant dispersion is observed due to grain boundary friction in the compositions. Moreover, the contact area is many times higher, and there is a high concentration of nonequilibrium defects and internal stresses. Dispersion causes new particle surfaces, while centers with increased activity are formed on the newly formed surfaces.
Upon activation, oxide layers break down and adsorb films on powder particles. Moreover, the following processes can be observed: crystal deformation; heat generation; local growth in temperature and pressure; shear stresses and changes in the size of microblocks; accumulation of crystallites; phase transformations; amorphization; breakdown of chemical bonds; acceleration of diffusion processes; “diffusion mixing”. Intense mechanical impact on the components increases reactivity due to plastic deformation of the crystalline structure, while the oxide is removed and adsorbs layers from the surface of the particles in the powder mixture. The use of HMT powder compositions makes it possible to create nanostructured composite coatings due to the increased reactivity in the HMT process, which provides clear technological advantages.
Figure 2 shows the dependence of the average particle size of the powder compositions GYYZO, material 1, and material 2 on the HMT time, obtained as a result of the statistical processing of the experimental data in the SPSS Statistica 10.0 environment.
The dependence of the average size of powder particles (D, μm) on activation time (T, min) is described by Equations (5)–(7), obtained by processing the experimental data in the Statistica 10.0 program.
Powder composition “GYYZO”:
D = 29.31 − 0.768·T + 0.0039·T2,
Powder composition “material 1”:
D = 59.46 − 1.084·T − 0.0073·T2,
Powder composition “material 2”:
D = 59.68 − 0.661·T − 0.018·T2.
Thus, it was established that during the HMT of the multicomponent powder materials, an increase in the HMT produced more finely dispersed powder compositions, homogenization was observed, and, due to hardening, the strength characteristics increased. New powder materials were formed, and they corresponded to the materials necessary to obtain layers of composites, which could be confirmed by the chemical analysis of the composites formed by VPS.
At the moment of loading into the drum for the HMT, the powder mixture, which had a given average concentration, was non-uniform in volume. The parameters and degree of concentration and heterogeneity of the components, along with the parameters of the dynamic effect of the grinding balls, were the determining factors in increasing the reactivity of the mixture being processed. The stress field and energy accumulation during the HMT significantly depended on the loading and size of the powder particles, as well as on the speed of the balls. At the initial stage of the process, if there were large particles in the powder compositions, high speeds were required to grind them. After reaching a certain particle size in powder composition, it was advisable to reduce the speed. During this process, plastic deformation was obtained, accompanied by a temperature increase, in addition to grinding.
The analysis showed that during the HMT, at the initial point the surface films were destructed and a slight increase in the specific surface area was observed. This destruction activated the reacting components of the powder mixtures. Beyond the plastic deformation, the HMT led to an increase in temperature at the point of contact with the grinding media. The increase in temperature caused the particles to be welded into conglomerates. The exterior and morphology of the surface in the initial powder compositions after the HMT are shown in Figure 3. When the HMT lasted for more than 5 min, the powder particles formed agglomerates, and their contact area increased many times. The number of grinding balls involved in the process influenced the formation of agglomerates. As the number of balls increased, the generated energy grew and nonequilibrium defects concentrated on the newly formed surface [60]. When the HMT lasted for more than 15–25 min, the newly formed agglomerates were destructed and then again participated in repeated mechanical fusion. This increased the defectiveness of the powder composition, which in turn increased the reaction activity. When the HMT lasted longer (25–35 min), the chemical composition of the particles was graded, the collision was stabilized, and the temperature increased, which also led to the welding of the particles.
The structure and properties of the composite layers are determined by contact processes during the impact, deformation, solidification, cooling, and interaction of the particles during movement in a gas flow. Since the layers are formed by the consistent deposition of many deformed particles, the formation of pores at the boundaries of the joints is absolutely inevitable. Obviously, the structure of the layers depends on the granulometric composition of the sprayed material. So, the finer the sprayed powder is, the less likely it is to form pores. Experience shows that the use of small powders for spraying causes some technological difficulties, such as the burning out of some components of the sprayed material in the gas flow. Consequently, optimization of the particle size distribution is one of the important stages in composite layer technology. To obtain high-quality layers, the optimal granulometric powder composition is 0.5–10 microns. When using powder of smaller fractions (less than 0.4 microns), the particles stick together and, as a result, the technological process is disrupted (the powder dispenser channel becomes clogged). The powder size before mechanical alloying was 30–40 µm, and after mechanical alloying, it was 0.8–10 µm.
As a result of the HMT, internal stresses arise in local powder microvolumes. The relaxation of these stresses depends on the properties of the material and the loading conditions. According to existing theories, mechanochemical transformations are initiated by the heat released during the powder processing, by dislocation energy during the plastic deformation, by the elastic energy release, and by numerous interphase boundaries. During the HMT, the crystal lattice of the metal is deformed. As a result, there is spatial and energetic inhomogeneity of the surface, which increases the defectiveness and energy saturation due to the plastic deformation and crushing [61,62,63]. When the plasma jet interacts with the HMT powders, the accumulated energy is released, which changes the material properties. As a result, the adhesion and cohesion of the layers increases and their porosity decreases. This can probably be explained by the formation of defects during the HMT process, the energy of which is released during VPS.

3.2. Structural Studies of Composites

Macro- and microanalysis of the composite layers obtained using the proven technology showed that their structure was quite dense. The interface between the layers had no visible defects (Figure 4a–d). When particles of the powder materials were passed through a high-temperature jet, they heated up and when they hit the substrate, they hardened in the form of deformed disks with a diameter of 1–8 microns and a thickness of 1–2 microns. The grain size in the layers of the composite materials Hastelloy X/Rene N6–GYYZO–material 1 and Hastelloy X/Rene N6–GYYZO–material 2, obtained by VPS, was as follows: GYYZO ex 180 to 470 nm (Figure 4e); material 1 ex 52 to 215 nm (Figure 4f,g); material 2 ex 45 to 218 nm (Figure 4g,h). The thickness of the layers of the hybrid composite materials Hastelloy X/Rene N6–GYYZO–material 1 and Hastelloy X/Rene N6–GYYZO–material 2 was as follows: GYYZO ex 0.1 to 0.2 mm; material 1 and material 2, ex 1 to 1.3 mm. This thickness of the composite layers has optimal mechanical properties for hybrid composites.
Figure 4a shows the structure of the Hastelloy X–GYYZO–material 1 composite. Figure 4b shows an enlarged part of the structure (interface between the layers of the GYYZO alloy and material 1) of the Hastelloy X/Rene N6–GYYZO–material 1 composite. Figure 4c demonstrates a scaled-up part of the composites (the interface between the GYYZO and material 2 alloy layers) of the Hastelloy X/Rene N6–GYYZO–material 2 composite.
The results of the X-ray phase analysis showed that the thermal barrier layer of GYYZO consisted of a tetragonal ZrO2 phase (JCPDS cards no. 37-1484 and 17-0923), a cubic Gd2O3 phase (JCPDS cards no. 12-0797 and 43-1014), a cubic Yb2O3 phase (JCPDS card no. 41-1106), and a cubic Y2O3 phase (JCPDS cards no. 71-0099, 44-0399 and 89-5592) (Figure 5a); layer material 1 (Table 1) consisted of a cubic Ni phase (ICCD: 04-0850; JCPDS cards no. 87-0712 and 01-078-7533), a cubic HfC phase (JCPDS cards no. 39-1491 and 65–7326), a tetragonal ZrO2 phase (JCPDS cards no. 37-1484 and 17-0923, 80-0965), a cubic Gd2O3 phase, a cubic Yb2O3 phase (JCPDS cards no. 43-1037and 41-1106), a cubic Y2O3 phase (JCPDS cards no. 71-0099, 44-0399 and 89-5592), a hexagonal Re phase (JCPDS card no. 05-0702), a cubic NiO phase (JCPDS cards no. 22-1189, 47-1049 and 04-0835) (Figure 5b); and layer material 2 (Table 1) consisted of a cubic Ni phase, a cubic HfC phase, a cubic TaC phase (JCPDS cards no. 061–8832, 77-0205 and 89-3831), a tetragonal ZrO2 phase, a cubic Gd2O3 phase, a cubic Yb2O3 phase, a cubic Y2O3 phase, a hexagonal Re phase, and a cubic NiO phase (Figure 5c).
During the subsequent heat treatment (annealing) for 2 h at a temperature of 900 °C, the degree of crystallinity of the ZrO2 and Gd2O3 phases increased. Thus, the nature of the diffraction pattern changed: the number of reflections related to ZrO2 and Gd2O3 sharply increased and became more intense. The maximum number of reflections was observed in the diffraction patterns of the samples subjected to thermal treatment at 900 °C. The monoclinic and tetragonal phases of ZrO2 and the monoclinic and cubic phases of Gd2O3 were present together in all the samples, and the ratio of these phases did not change depending on the formation temperature and subsequent heat treatment.
The monoclinic phase of ZrO2 had the unit cell parameters a = 5.1477 × 10−10 m, b = 5.203 × 10−10 m, c = 5.3156 × 10−10 m, and β = 80.77; for the tetragonal phase ZrO2, a = 5.12 × 10−10 m and b = 5.25 × 10−10 m. The ordering of the initial phase and the size of its grains have a great influence on the transition interval. In this case, the high dispersion of the oxide component contributed to the stabilization of the tetragonal form of ZrO2 and to an increase in the transition interval of the monoclinic ZrO2 to the tetragonal one. In addition, during a phase transformation, an activated state may arise due to the loosening of the lattice and the presence of stresses that are inevitable during the restructuring of the structure.

3.3. Properties of Surface-Modified Layers of Hybrid Composites

3.3.1. Porosity

One of the important characteristics that largely determines the performance of porous materials is the pore size and their percentage [64]. Maximum pore sizes are an important characteristic of materials used for porous filters, since they determine the maximum size of the pollutant particles that can theoretically pass through the filter. Average pore sizes are usually used as a characteristic for their comparison and as a determining size when processing experimental data on the hydraulic resistance of a porous medium in criterion form. The metallographic observations of the material structures under study indicated a fairly high uniformity of pore sizes in the volume of the composite layers. It is known that a porous material has a lower hardness than a pore-free material [65,66]. Therefore, it is relevant to study the effect of the refractory ceramic content and oxide phases on the material hardness.
Figure 6 shows how the porosity (Figure 6a) and the average pore size (Figure 6b) depended on the content of the ceramic phase in the composite layers. By analyzing the obtained dependences, one can note an increase in the porosity and pore size with an increasing concentration of the ceramic phase.
Some differences in the average pore sizes were observed in the “material 1” and “material 2” layers. This is explained by the large melting temperature gradient of the main ceramic phases of HfC, TaC, Gd2O3, ZrO2, and the nickel matrix.
As a result of the experimental and mathematical processing (Statistica 10.0) of our research data, the dependencies were constructed and empirical mathematical equations were compiled for the influence of the refractory phase content on the porosity (P, %) (Equation (8)) and the average pore size (D, µm) (Equation (9)):
P = 0.3843 + 0.1778·n − 0.0006·n2
D = 0.1763 + 0.0068·n − 3.3107·10−5·n2
where n—refractory phase content, %.
It should also be noted that the presented results are a fairly convincing illustration of the possibility to precisely control the volume and pore size in composite layers from almost 0.5 to 25% by changing the technological parameters of layer formation and the refractory phase content of HfC and TaC.

3.3.2. Microhardness

The microhardness of the hybrid composite materials Hastelloy X–GYYZO–material 1 and Hastelloy X–GYYZO–material 2 was as follows: super alloy Hastelloy X, HV0.2 = 3.8–3.95 GPa; layer GYYZO, HV0.3 = 16.1–16.7 GPa; layer material 1 (Table 1), HV0.3 =18.3–18.8 GPa; layer material 2 (Table 1), HV0.3 =19.1–19.6 GPa. These data suggest that the “layered structures” increased the mechanical characteristics of the material, in particular, they increased the wear resistance and improved the strength characteristics.

3.3.3. Mechanical Properties

The mechanical properties of the composites were determined by bend testing. The transition from a dispersion-strengthened to a porous structure was very clearly recorded when testing the composite samples for bending. Figure 7 shows the dependence of the change in the flexural strength of the hybrid composite materials Hastelloy X–GYYZO–material 1 and Hastelloy X–GYYZO–material 2 depending on the mass fraction of the ceramic phase content (in material 1 and material 2 (Table 1)). On the graph (Figure 7), the transition from the dispersion-strengthened zone to the porous one is clearly visible. In the transition zone, the bending strength σu reaches its maximum and decreases with increasing porosity. In all cases, during testing, there was a brittle fracture of the porous surface-modified layers of the composites.
As a result of the mathematical processing (Statistica 10.0) of the experimental data, dependencies were constructed and empirical mathematical equations were compiled for the influence of the refractory carbide phase content on the tensile strength (σu) on bending (Equation (10)):
σU = 675.9046 + 1.0731·nc − 0.0157·nc2, (MPa)
where nc—content of carbide refractory phases in the material, %.
Strengthening the ceramic phases in dispersion-strengthened materials and simultaneously strengthening the metal matrix reduces their ductility. The energy transferred to the body during its deformation can be dissipated through plastic flow or accumulate in the vicinity of the stress concentrators. These can be inclusions and cavities (pores) that have interfaces with the matrix material. Plastic flow during deformation causes an additional increase in stored energy in the area of the concentrators, as a result of which cracks may appear in the material. The problem of taking into account plastic flow around inclusions and cavities (pores) and the formation of cracks in this area is quite complex. The difficulty lies in the need to take into account not only the size of the inclusions and the properties of the matrix materials and ceramic phases but also the characteristics of the interphase interaction at the interface between these inclusions and the matrix material.

3.3.4. Heat Resistance

The heat resistance of composite materials is one of the most important characteristics that determine their use for creating thermal protective coatings. For porous materials produced by powder metallurgy methods, the higher the density of the oxidizing gases, the larger the average pore size, and the higher the concentration of oxidizing gas, the higher the intensity of oxidation.
Figure 8 shows the kinetic curves of the surface layer oxidation (material 1; material 2; Ni-Al2O3 (24% wt.); NiCrY-Al2O3 (10% wt.); NiCrAl-ZrO2 (37% wt.); Hastelloy X; Rene N6), upon heating in an air atmosphere at a temperature of 1250 °C. The main increase in mass in the samples occurred during the first hours; subsequently, this dependence was almost linear. In one hour of testing, the increase in scale mass in the “material 1” was 0.37 g/m2; in 5 h—2.15 g/m2, and after 25 h—7.1 g/m2.
Analyzing the obtained dependences, it is easy to see that the developed material 1 and material 2 are characterized by the greatest resistance to oxidation (Figure 8).
As a result of the mathematical processing (Statistica 10.0) of the experimental data, dependencies were constructed and empirical mathematical equations were compiled for the influence of the test time (t, hour) on the mass gain (∆m, g/m2) (Equations (11)–(15)):
Material 1: ∆m = 2.0715 + 0.2067·t − 0.0005·t2;
Material 2: ∆m = 2.5318 + 0.1938·t + 0.0005·t2;
Ni-Al2O3: ∆m = 5.5517 + 2.0203·t − 0.0212·t2;
NiCrAl-ZrO2: ∆m = 31.6833 + 4.0202·t − 0.0368·t2;
NiCrY-Al2O3: ∆m = 50.7667 + 3.2858·t − 0.0573·t2;
Hastelloy X: ∆m = 108.017 − 1.33·t + 0.3674·t2;
Rene N6: ∆m = 53.817 + 4.174·t − 0.0205·t2.
The introduction of dispersed oxides significantly increased the heat resistance of the alloys due to the formation of a surface film consisting of the oxides Gd2O3, Al2O3, and Y2O3. The refractory particles of the oxide ceramic phases, being the centers of growth for a thin protective layer firmly adhered to the base, inhibit the oxidation of porous metal ceramics. A change in the morphology of oxide films in a more favorable direction is achieved by microalloying with gadolinium, yttrium, and aluminum oxides.

3.3.5. Wear Resistance

The composite layers were tested for wear during dry friction against a hard rotating disk on a 2070 SMT-1(Moscow, Russia) testing machine with a disk rotation speed v = 0.5–2 m·s−1 and a pressure P = 2–12 MPa. The temperature was recorded in a contact zone. The wear rate was assessed according to the experimental data using the Statistica v10.0 package in the SPSS environment. Figure 9 shows the results of the wear test of the samples made of the Hastelloy X alloy with composite coatings with a thickness of 0.9–1.5 mm, obtained by VPS, according to the disc-to-disc scheme.
Figure 9a shows the results of the friction-wear tests of the composite materials NiCrY-Al2O3, NiCrAl-ZrO2, GYYZO + material 1, GYYZO + material 2, and Hastelloy X (NiCrFeMo). The thickness of the composite layers was 1.5 mm, and the samples were tested on the friction-wear testing machine 2070 CMT-1 (Figure 9b). The tests were carried out in dry friction with a disk sliding speed of 2 m/s. The temperature at the contact point was measured during loading. The data were processed with the Statistica v10.0 package in the SPSS environment. The analysis of the results shows (Figure 9a) that the material Hastelloy X (NiCrFeMo) (I = 46.16 × 10−6 g/m, disk pressure P = 9.4 MPa) had the highest wear rate (Figure 9a, curve 1), and the lowest wear rate belonged to the composite material GYYZO + material 2 (I = 28.12 × 10−6 g/m, disk pressure P = 9.4 MPa) (Figure 9a, curve 1).
An increased wear resistance in the composite layers of material 1 and material 2 was achieved by reinforcing the layers with superhard carbide inclusions of HfC and TaC.
The wear rate (I) from the pressure (P) of the disk changes according to a polynomial dependence (Figure 9a) can be described by the empirical Equations (18)–(22) obtained from the mathematical analysis of the experimental data in the Statistica v10.0 software:
GYYZO + material 1: I = −8.4859 + 3.1269·P + 0.0816·P2;
GYYZO + material 2: I = −8.2749 + 2.9432·P + 0.0776·P2;
Hastelloy X (NiCrFeMo): I = −10.0661 + 4.4004·P + 0.1676·P2;
NiCrAl-ZrO2 (37% wt.): I = −6.9886 + 2.9841·P + 0.1961·P2;
NiCrY-Al2O3 (10% wt.): I = −6.7152 + 2.9324·P + 0.2517·P2.
The positive test results for the wear and heat resistance give reason to recommend the composite materials Hastelloy X/Rene N6 + GYYZO + material 1 and Hastelloy X/Rene N6+ GYYZO + material 2 to improve the reliability of various structural elements operating under conditions of high temperature and frictional –cyclic loading.

4. Conclusions

Hybrid layered reinforced materials can successfully ensure the reliability and durability and expand the functionality of the high-temperature components and assemblies of supercritical and ultra-supercritical power plants and oil, gas, and petrochemical equipment operating under multifactorial influences (temperature, force, deformation). As a result of this research, surface reinforced ceramic composite materials with a gradient distribution of properties have been developed. These properties included thermal barrier layers and layers of material with carbide and oxide inclusions, with high heat resistance, strength, and wear resistance, providing the specified functional and mechanical properties.
This research has justified the choice of the composition of the thermal barrier (Gd2O3-Yb2O3-Y2O3-ZrO2) and the functional layers (Ni-based layers reinforced with refractory carbide particles of Hf, Ta, oxide particles of Gd, Al, Zr, etc.), as well as the heat-resistant and wear-resistant composites used as functional protective materials in the enterprises of the energy and oil and gas industries. The choice of the equipment and technology for the formation of the layered materials has also been justified. As a formation technology, HMT (mechanical alloying) of the powder compositions was proposed, followed by plasma spraying in a vacuum. As a result of the conducted research, it was established that the use of high-energy mechanical processing of the materials before plasma spraying in a vacuum leads to an increase in adhesion and cohesion of the composite layers and a decrease in their porosity. The adhesive strength of the layered composite materials was Hastelloy X/GYYZO—32 MPa; GYYZO/material 1—34 MPa; Rene N6/GYYZO—29 MPa; GYYZO/material 2—31 MPa. The adhesive strength of the materials after HMT was Hastelloy X/GYYZO—39 MPa; GYYZO/material 1—43 MPa; Rene N6/GYYZO—34 MPa; GYYZO/material 2—38 MPa. The structure and phase composition of the composites were studied. The grain size in the layers of the composite materials was Hastelloy X/Rene N6–GYYZO–material 1, Hastelloy X/Rene N6–GYYZO–material 2, and the results obtained by VPS were as follows: GYYZO—180–470 nm; material 1—52–215 nm; material 2—45–218 nm. The microhardness of the alloy layers of the hybrid composite materials Hastelloy X–GYYZO–material 1 and Hastelloy X–GYYZO–material 2 was super alloy Hastelloy X, HV0.2 = 3.8–3.95 GPa; layer GYYZO, HV0.3 = 16.1–16.7 GPa; layer material 1, HV0.3 =18.3–18.8 GPa; layer material 2, HV0.3 =19.1–19.6 GPa. The influence of the refractory phases of HfC and TaC on the strength of the composites was studied. It was found that the maximum strength (710–715 MPa) in the composites Hastelloy X–GYYZO–material 1 and Hastelloy X–GYYZO–material 2 is achieved with a content of HfC—TaC—27–28%. The heat resistance of the composites was studied: material 1 had the highest resistance to high-temperature oxidation, while Hastelloy X had the lowest resistance. The heat resistance of the composites was compared (GYYZO + material 1, GYYZO + material 2) with previously known materials (Ni-Al2O3 (24% wt.); NiCrY-Al2O3 (10% wt.); NiCrAl-ZrO2 (37% wt.)). The wear resistance of the composites was studied: material Hastelloy X (NiCrFeMo) (I = 46.16 × 10−6 g/m, disk pressure P = 9.4 MPa) had the highest wear rate, and the lowest wear rate belonged to the composite material GYYZO + material 2 (I = 28.12 × 10−6 g/m, disk pressure P = 9.4 MPa).
The test results show the developed composites are promising for use in various industries, including the energy industry (components and assemblies of thermal power plants) and oil and gas industry (turbine blades, etc.), where high heat and wear resistance of materials is required.

Author Contributions

Conceptualization, P.R., C.Z. and P.S.; methodology, P.R., C.Z., P.S. and A.R.; software, P.R., C.Z., P.S., G.K. and M.S.; validation, P.R., C.Z., P.S., A.R. and G.K.; formal analysis, P.R., P.S. and A.R.; investigation, P.R., C.Z., A.R., G.K. and M.S.; resources, P.R.; data curation, P.R., P.S. and G.K.; writing—original draft, P.R.; supervision, P.S.; project administration, P.R.; funding acquisition, P.R. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by Russian Science Foundation grant number 23-23-00074.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data are contained within the article.

Acknowledgments

This work was carried out under the financial support of the Russian Science Foundation No. 23-23-00074.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Multifunctional technological complex for the formation of hybrid composites in a protective environment—(a); plasma spraying of oxide layers—(b).
Figure 1. Multifunctional technological complex for the formation of hybrid composites in a protective environment—(a); plasma spraying of oxide layers—(b).
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Figure 2. Effect of HMT time on powder particle size: GYYZO—(a); material 1, material 2—(b).
Figure 2. Effect of HMT time on powder particle size: GYYZO—(a); material 1, material 2—(b).
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Figure 3. Powder compositions, crushed and mechanically activated in a planetary mill for 30 min: GYYZO—(a,b); material 1—(c,d); material 2—(e,f).
Figure 3. Powder compositions, crushed and mechanically activated in a planetary mill for 30 min: GYYZO—(a,b); material 1—(c,d); material 2—(e,f).
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Figure 4. Structure of layered composite materials: Hastelloy X–GYYZO–material 1 (cross section of samples)—(a); structure of GYYZO layer–material 1 layer (cross section of samples)—(b); structure of GYYZO layer–material 2 layer (cross section of samples)—(c); GYYZO layer structure—(d); material 1 layer structure—(e,f); material 2 layer structure—(g,h).
Figure 4. Structure of layered composite materials: Hastelloy X–GYYZO–material 1 (cross section of samples)—(a); structure of GYYZO layer–material 1 layer (cross section of samples)—(b); structure of GYYZO layer–material 2 layer (cross section of samples)—(c); GYYZO layer structure—(d); material 1 layer structure—(e,f); material 2 layer structure—(g,h).
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Figure 5. X-ray diffraction patterns of hybrid composite materials: GYYZO—(a); material 1—(b); material 2—(c).
Figure 5. X-ray diffraction patterns of hybrid composite materials: GYYZO—(a); material 1—(b); material 2—(c).
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Figure 6. Influence of the refractory ceramic phase content on the porosity—(a) and average pore size—(b) of the composite layers.
Figure 6. Influence of the refractory ceramic phase content on the porosity—(a) and average pore size—(b) of the composite layers.
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Figure 7. Influence of the carbide phase content (HfC, TaC) in composite layers on the flexural strength.
Figure 7. Influence of the carbide phase content (HfC, TaC) in composite layers on the flexural strength.
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Figure 8. Dependence of mass gain (∆m) on exposure time (t) at temperature T = 1250 °C: 1—material 1; 2—material 2; 3—Ni-Al2O3 (24% wt.); 4—NiCrY-Al2O3 (10% wt.); 5—NiCrAl-ZrO2 (37% wt.); 6—Rene N6; 7—Hastelloy X.
Figure 8. Dependence of mass gain (∆m) on exposure time (t) at temperature T = 1250 °C: 1—material 1; 2—material 2; 3—Ni-Al2O3 (24% wt.); 4—NiCrY-Al2O3 (10% wt.); 5—NiCrAl-ZrO2 (37% wt.); 6—Rene N6; 7—Hastelloy X.
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Figure 9. The dependence of the wear intensity I on the pressure of the disk P and the sliding speed of the disk: Hastelloy X (NiCrFeMo)—1; NiCrY-Al2O3 (10% wt.)—2; NiCrAl-ZrO2 (37% wt.)—3; GYYZO + material 1—4; GYYZO + material 2—5; GYYZO + material 2—(a); photo of the friction-wear test at a disk sliding speed of 2 m/s—(b).
Figure 9. The dependence of the wear intensity I on the pressure of the disk P and the sliding speed of the disk: Hastelloy X (NiCrFeMo)—1; NiCrY-Al2O3 (10% wt.)—2; NiCrAl-ZrO2 (37% wt.)—3; GYYZO + material 1—4; GYYZO + material 2—5; GYYZO + material 2—(a); photo of the friction-wear test at a disk sliding speed of 2 m/s—(b).
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Table 1. Contents of basic chemical elements, wt.%.
Table 1. Contents of basic chemical elements, wt.%.
MaterialNiMoCrFeWCoMnHfCReSiTaCZrO2Y2O3CuGd2O3Yb2O3TiAlAl2O3TaHf
Hastelloy X43.39.521.820.10.71.90.8--0.9---0.4--0.20.4---
Rene N657.41.44.2-612.5--5.4--------5.75-7.20.15
Gd2O3-Yb2O3-Y2O3-ZrO2 (GYYZO)-----------535-357-----
Material 130------204--87-96--16--
Material 230------163-1265-85--15--
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Rusinov, P.; Zhang, C.; Sereda, P.; Rusinova, A.; Kurapov, G.; Semadeni, M. Development and Research of New Hybrid Composites with Increased Requirements for Heat and Wear Resistance. Ceramics 2025, 8, 8. https://doi.org/10.3390/ceramics8010008

AMA Style

Rusinov P, Zhang C, Sereda P, Rusinova A, Kurapov G, Semadeni M. Development and Research of New Hybrid Composites with Increased Requirements for Heat and Wear Resistance. Ceramics. 2025; 8(1):8. https://doi.org/10.3390/ceramics8010008

Chicago/Turabian Style

Rusinov, Peter, Chao Zhang, Polina Sereda, Anastasia Rusinova, George Kurapov, and Maxim Semadeni. 2025. "Development and Research of New Hybrid Composites with Increased Requirements for Heat and Wear Resistance" Ceramics 8, no. 1: 8. https://doi.org/10.3390/ceramics8010008

APA Style

Rusinov, P., Zhang, C., Sereda, P., Rusinova, A., Kurapov, G., & Semadeni, M. (2025). Development and Research of New Hybrid Composites with Increased Requirements for Heat and Wear Resistance. Ceramics, 8(1), 8. https://doi.org/10.3390/ceramics8010008

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