High-temperature and high-pressure nanometer high-conductivity heat exchange material

Abstract

This invention relates to the field of heat exchange materials technology, and discloses a high-temperature, high-pressure nano-conductivity heat exchange material. The heat exchange material comprises a metal matrix, a linearly distributed metal-ceramic gradient transition layer, and a nano-composite ceramic surface layer. The surface layer consists of 90-95 parts of nano-alumina and zirconium oxide composite powder and 5-10 parts of cubic phase nano-titanium nitride powder. The former is fused and stacked to form a continuous ceramic matrix, while the latter is monodispersed and anchored within it, with adjacent particles in contact. The heat exchange material is prepared by gradient layer spraying with a linearly decreasing plasma current as the ceramic composition increases, followed by back-side cooling of the matrix. The surface layer is sprayed using a coaxial air knife in a reducing atmosphere combined with a low-enthalpy jet to inhibit oxidation and vacuum diffusion annealing. This invention synergistically regulates the above-mentioned structure and interface state to achieve a balance between high thermal conductivity, excellent corrosion resistance, and high bonding strength, which helps extend the service life of heat exchange components under high-temperature and high-pressure environments.

Description

Technical Field

[0001] This invention relates to the field of heat exchange materials technology, and in particular to a high-temperature, high-pressure nano-conductivity heat exchange material. Background Technology

[0002] Heat exchange materials are the core medium in industrial heat energy conversion and transfer systems. The physicochemical properties of heat exchange materials determine the energy utilization efficiency and operational safety of heat exchange equipment. Under high temperature, high pressure, and highly corrosive conditions, a protective coating is usually prepared on the surface of the metal substrate. The protective coating is used to protect the metal substrate from corrosion and wear, while maintaining effective heat transfer to ensure the stable operation of the heat exchange equipment in harsh environments.

[0003] Existing technologies typically employ thermal spraying processes to prepare oxide ceramic coatings or metal-ceramic composite coatings on the surface of metal workpieces. Oxide ceramic materials such as alumina and zirconium oxide are often used to isolate corrosive media due to their excellent chemical stability and high-temperature resistance. Metal-ceramic composite coatings attempt to improve the toughness of the coating and its bonding performance with the metal substrate by adding metal powder to the ceramic matrix. Existing technical solutions mainly focus on extending the service life of heat exchange equipment through physical isolation under normal operating conditions.

[0004] However, due to the inherently low thermal conductivity of oxide ceramic materials and the significant difference in thermal expansion coefficients between the ceramic coating and the metal substrate, existing technologies suffer from performance defects in high-temperature and high-pressure environments. Thicker ceramic coatings significantly increase thermal resistance, leading to reduced heat transfer efficiency in heat exchange equipment. While reducing the thickness of the ceramic coating can improve heat transfer, it is difficult to effectively prevent the penetration of corrosive media. Simultaneously, the mismatch in thermal expansion coefficients causes accumulated thermal stress at the heterogeneous interface during alternating hot and cold thermal cycles, resulting in coating cracking or peeling. Furthermore, existing technologies that attempt to introduce high thermal conductivity phases to improve thermal conductivity often suffer from corrosion resistance due to oxidation or agglomeration of the thermally conductive phase during high-temperature spraying, which disrupts the continuous and dense structure of the ceramic layer. This makes it difficult for existing technologies to simultaneously meet the comprehensive requirements of high thermal conductivity, high corrosion resistance, and high bonding strength. Summary of the Invention

[0005] The purpose of this invention is to provide a high-temperature and high-pressure nano-conductivity heat exchange material, which solves the problems of existing metal-ceramic heat exchange coatings being unable to simultaneously achieve high thermal conductivity and excellent corrosion resistance under high-temperature and high-pressure conditions, and the coatings easily peeling off and failing due to mismatched coefficients of thermal expansion.

[0006] The first aspect of this invention provides a high-temperature, high-pressure nano-conductivity heat exchange material, comprising a metal matrix, a metal-ceramic gradient transition layer formed on the surface of the metal matrix, and a nanocomposite ceramic surface layer located on the metal-ceramic gradient transition layer; the nanocomposite ceramic surface layer is prepared from raw materials comprising the following parts by weight:

[0007] The nanostructured alumina and zirconium oxide composite powder comprises 90 to 95 parts; cubic phase nano-titanium nitride powder comprises 5 to 10 parts; the nanostructured alumina and zirconium oxide composite powder is melt-stacking to form a continuous ceramic matrix, and the cubic phase nano-titanium nitride powder is anchored in the continuous ceramic matrix in a monodisperse form, with adjacent nano-titanium nitride particles in contact with each other.

[0008] By employing the above technical solution, the high-temperature, high-pressure nano-high-conductivity heat exchange material utilizes the high intrinsic thermal conductivity of cubic phase titanium nitride nanoparticles, combined with the specific dispersion morphology of cubic phase titanium nitride nanoparticles in an alumina and zirconium oxide composite matrix, to achieve a nonlinear improvement in thermal performance. The working mechanism of the high-temperature, high-pressure nano-high-conductivity heat exchange material is as follows:

[0009] First, the thermal conductivity enhancement mechanism. When the amount of high thermal conductivity titanium nitride powder added is controlled within a specific range of 5 to 10 parts, monodisperse titanium nitride nanoparticles can be uniformly distributed in the ceramic matrix, and adjacent particles can make contact with each other. These titanium nitride particles are distributed at the grain boundaries and within the ceramic matrix, and the contact between adjacent particles forms a path for efficient heat transfer, effectively reducing the hindering effect of alumina and zirconium oxide grain boundaries on heat transfer, thereby significantly improving the overall thermal conductivity of the high-temperature and high-pressure nano-high-conductivity heat exchange material.

[0010] Second, the corrosion channel blocking mechanism. By strictly limiting the titanium nitride content to below 10 parts and ensuring it is monodisperse, excessive agglomeration of conductive phase particles to prevent the formation of a macroscopically interconnected conductive network is avoided. The alumina and zirconium oxide composite matrix remains as a continuous phase surrounding the titanium nitride, maintaining the density of the physical barrier layer and cutting off the electrochemical pathway for chloride ion corrosion media to penetrate inward. This improves thermal conductivity without sacrificing corrosion resistance.

[0011] Third, stress regulation mechanism. As a functionally graded material, the metal-ceramic gradient transition layer alleviates the abrupt change in the coefficient of thermal expansion between the metal matrix and the nanocomposite ceramic surface through continuous changes in composition, reduces the interfacial shear stress during thermal shock, and prevents coating peeling.

[0012] Preferably, the nanostructured alumina and zirconium oxide composite powder is a spherical agglomerate that has undergone spray granulation and sintering treatment, with a particle size range of 5 to 45 micrometers; the cubic phase nano-titanium nitride powder has an average particle size of 40 to 60 nanometers.

[0013] By adopting the above technical solution, the micron-sized spherical aggregates ensure the fluidity and deposition efficiency during spraying, while the nano-sized titanium nitride particles fill the pores and surface defects formed by the accumulation of micron-sized particles, thereby improving the density of the nanocomposite ceramic surface layer.

[0014] Preferably, the metal-ceramic gradient transition layer is formed by melt deposition of nickel-chromium alloy powder and alumina and titanium oxide composite powder, wherein the alumina and titanium oxide composite powder is the ceramic component of the metal-ceramic gradient transition layer.

[0015] By adopting the above technical solution, the nickel-chromium alloy provides metallurgical affinity with the metal matrix and high-temperature oxidation resistance. The introduction of titanium oxide improves the wettability of the ceramic components and enhances the interfacial bonding between the metal phase and the ceramic components.

[0016] Preferably, in the metal-ceramic gradient transition layer, from the side closer to the metal substrate to the side closer to the nanocomposite ceramic surface, the mass percentage of nickel-chromium alloy powder linearly decreases from 100% to 20%-50%, while the mass percentage of alumina and titanium oxide composite powder linearly increases from 0% to 50%-80%.

[0017] By adopting the above technical solution, the composition of the transition layer gradually changes from the side closer to the metal substrate to the side closer to the surface of the nanocomposite ceramic, which avoids stress concentration caused by sudden changes in composition, improves the ability of high-temperature and high-pressure nano-conductivity heat exchange materials to resist alternating hot and cold temperatures, and extends their fatigue service life.

[0018] Preferably, the metal matrix and the metal-ceramic gradient transition layer, as well as the components within the metal-ceramic gradient transition layer, are all interatomic diffused and bonded through vacuum high-temperature diffusion; the oxygen content of the surface layer of the nanocomposite ceramic is 0.3-0.6 wt%.

[0019] By adopting the above technical solution, the inter-atomic diffusion bonding transforms the original mechanical interlocking interface between the metal matrix and the metal-ceramic gradient transition layer, and between the components inside the transition layer, into an atomic-level bonding state, thereby improving the bonding strength of each interface. The low oxygen content of 0.3-0.6 wt% on the surface of the nanocomposite ceramic confirms that titanium nitride has not undergone oxidation and deterioration, and can maintain the cubic crystal structure and high thermal conductivity of titanium nitride.

[0020] Preferably, the metal substrate is a carbon steel substrate, and the surface roughness Ra of the metal substrate is 4.5 to 6.0 micrometers; the thickness of the metal-ceramic gradient transition layer is 50 to 150 micrometers, and the thickness of the nanocomposite ceramic surface layer is 30 to 80 micrometers.

[0021] By adopting the above technical solutions, the specific roughness provides sufficient mechanical anchoring points, and the reasonable thickness ratio provides sufficient corrosion resistance and stress buffer margin while ensuring controllable thermal resistance.

[0022] A second aspect of this invention provides a method for preparing a high-temperature, high-pressure nano-conductivity heat exchange material, comprising the following steps:

[0023] Provide a metal substrate and perform degreasing and sandblasting roughening treatment on the surface of the metal substrate;

[0024] An atmospheric plasma spraying process is used to spray a metal-ceramic gradient transition layer onto the surface of the metal substrate. During the spraying process, the plasma current is linearly reduced as the ceramic component content increases, and cooling is activated for the back of the metal substrate.

[0025] A nanocomposite ceramic surface layer is sprayed onto a metal-ceramic gradient transition layer. A coaxial annular air knife is applied at the spray gun outlet and nitrogen is introduced to create a local reducing atmosphere. Combined with plasma jet parameters of low enthalpy and high flow rate, the oxidation of cubic phase nano-titanium nitride powder is suppressed.

[0026] The coated workpiece is placed in a vacuum furnace for high-temperature diffusion annealing and then cooled in the furnace to obtain a high-temperature, high-pressure nano-conductivity heat exchange material.

[0027] By employing the above technical solution, the preparation method of high-temperature and high-pressure nano-conductivity heat transfer materials solves the problems of thermal stress accumulation and nano-phase oxidation during the deposition of heterogeneous materials through multi-dimensional process control. The specific process mechanism is as follows:

[0028] First, the thermal management mechanism of gradient layer spraying. During the deposition of the cermet gradient transition layer, as the ceramic content increases, the thermal conductivity of the cermet gradient transition layer decreases and its brittleness increases. At this time, linearly reducing the plasma current reduces the input heat, while activating back cooling to forcibly remove heat from the substrate can suppress the formation of tensile stress during coating solidification shrinkage and prevent the initiation of microcracks due to heat accumulation.

[0029] Secondly, the mechanism of local atmosphere protection and kinetic control. In an atmospheric environment, nano-titanium nitride readily undergoes an oxidation reaction with oxygen to produce titanium dioxide and nitrogen gas. The generated titanium dioxide has a much lower thermal conductivity than titanium nitride, which can lead to the breakage of the heat conduction network. The preparation method of high-temperature and high-pressure nano-high-conductivity heat exchange materials uses a coaxial annular air knife to create a local oxygen-deficient environment by passing nitrogen gas through it. Combined with low enthalpy and high flow rate parameters, this shortens the residence time of particles in the high-temperature flame, kinetically inhibiting the oxidation reaction and ensuring that titanium nitride is deposited in its original phase.

[0030] Finally, the diffusion annealing mechanism. Vacuum high-temperature treatment eliminates the interlayer porosity in the sprayed state, drives the diffusion of metal atoms and ceramic atoms across the interface to form a dense diffusion layer, and releases the residual thermal stress from the preparation process.

[0031] Preferably, during the process of spraying the metal-ceramic gradient transition layer, the plasma current is linearly reduced from 550 to 600 amperes to 450 to 500 amperes, and the surface temperature of the metal substrate is reduced from 160 to 200 degrees Celsius to 85 to 110 degrees Celsius.

[0032] By adopting the above technical solution, the precise current and temperature window ensures that the metal phase is fully melted and spread, while avoiding overheating and decomposition of the ceramic components or the generation of excessive solidification shrinkage stress.

[0033] Preferably, during the process of spraying the nanocomposite ceramic surface layer, the nitrogen flow rate introduced by the coaxial annular air knife is controlled at 35 to 50 liters per minute; the low enthalpy and high flow rate plasma jet parameters include: plasma current of 480 to 520 amperes, main gas argon flow rate of 52 to 60 liters per minute, and auxiliary gas hydrogen flow rate of 2.5 to 3.5 liters per minute.

[0034] By adopting the above technical solution, the high flow rate of argon gas provides strong momentum transfer, giving the nano-aggregates extremely high flight speed; an appropriate amount of hydrogen gas increases the jet enthalpy without causing overheating; and a sufficient flow of nitrogen gas effectively eliminates the surrounding air entrained into the plasma jet.

[0035] Preferably, the process parameters for high-temperature diffusion annealing are: a vacuum degree of 1.0 x 10⁻⁴ Pascals to 2.0 x 10⁻³ Pascals, a heating rate of 5 to 8 degrees Celsius per minute, an annealing temperature of 850 to 920 degrees Celsius, and a holding time of 2.5 to 4 hours.

[0036] By adopting the above technical solution, the annealing temperature and vacuum degree are exactly in the process window where iron, nickel and chromium elements diffuse actively but the grains do not grow severely, thus maximizing the bonding strength and minimizing the damage to the matrix structure.

[0037] In summary, the present invention has at least one of the following beneficial technical effects:

[0038] 1. This invention introduces a specific proportion of monodisperse cubic phase nano-titanium nitride powder into an alumina and zirconium oxide composite matrix. This allows the nano-titanium nitride powder to be anchored in a continuous ceramic matrix in a monodisperse form, with adjacent particles in contact with each other. The high intrinsic thermal conductivity of titanium nitride enables efficient heat transfer, which is beneficial to improving the overall heat transfer efficiency of the material. At the same time, the monodisperse distribution maintains the physical continuity of the insulating ceramic matrix, which helps to reduce the corrosion medium penetration channels formed by excessive agglomeration of the conductive phase. This is beneficial to achieving both high thermal conductivity and resistance to electrochemical corrosion.

[0039] 2. This invention employs a metal-ceramic gradient transition layer combined with a vacuum high-temperature diffusion annealing process, which alleviates the problem of mismatched thermal expansion coefficients between ceramic and metal heterogeneous materials. The linearly distributed composition gradient smooths the thermal stress distribution along the thickness direction, reducing the risk of stress concentration. The vacuum high-temperature treatment promotes atomic interdiffusion at the interface between the transition layer and the substrate, realizing interatomic diffusion bonding and transforming the original mechanical interlocking interface into an atomic-level bonding interface. This helps to improve the bonding strength and thermal shock resistance of the coating, and has a positive effect on the service life of the material under high temperature and high pressure complex working conditions.

[0040] 3. In the preparation process, this invention constructs a local reducing atmosphere using a coaxial annular air knife and combines it with plasma jet parameters of low enthalpy and high flow rate to kinetically suppress the oxidation behavior of nano-titanium nitride during high-temperature spraying. This process control is beneficial for retaining the original cubic crystal phase structure of titanium nitride in the coating, reducing the risk of thermal conductivity network breakage caused by oxidation to form low thermal conductivity titanium dioxide, and supporting the realization of the material's designed thermal properties. Attached Figure Description

[0041] Figure 1 This is a schematic diagram illustrating the nonlinear effect of the amount of nano-titanium nitride added in this invention on the thermal conductivity of the coating.

[0042] Figure 2 This invention provides a comparison of the electrodynamic polarization curves of the coating under simulated marine conditions.

[0043] Figure 3 This is a bar chart comparing the interfacial bonding strength of the coatings of this invention. Detailed Implementation

[0044] The technical solutions in the embodiments of the present invention will be clearly and completely described below with reference to the preparation examples, comparative examples, and test examples. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0045] Preparation Examples 1-3:

[0046] Preparation Example 1:

[0047] This preparation example provides a method for preparing a nanocomposite spraying powder with a TiN mass fraction of 5%, including the following steps:

[0048] 950g of Al2O3-ZrO2 nanostructure powder with a particle size of less than 45μm and greater than 5μm, which has been treated by spray granulation and sintering, and 50g of cubic crystalline phase nano-TiN powder with an average particle size of 50nm were selected.

[0049] The two powders were placed in the agate ball mill jar of a planetary ball mill, and zirconia grinding balls with diameters of 5 mm and 10 mm were added. The ball-to-material mass ratio was controlled at 5:1, and 1500 ml of anhydrous ethanol was added as a process control agent.

[0050] The ball mill was set to rotate at 150 r / min and low-energy ball milling was performed for 2 hours to ensure that the nano-TiN particles were uniformly attached to the pores and surface of the micron-sized Al2O3-ZrO2 agglomerates, while avoiding damage to the sphericity of the Al2O3-ZrO2 agglomerates.

[0051] The slurry after ball milling was removed and the grinding balls were filtered out. It was then placed in a rotary evaporator and dried by rotary evaporation under a vacuum of 0.08 MPa and a water bath temperature of 60°C until the solvent was completely evaporated and loose powder was obtained.

[0052] The dried powder was sieved through a 200-mesh standard sieve to remove large particle agglomerates, resulting in a TiN-doped composite spray powder with good flowability and a light gray color.

[0053] Preparation Example 2:

[0054] This preparation example provides a method for preparing a nanocomposite spraying powder with a TiN mass fraction of 8%, including the following steps:

[0055] 920g of Al2O3-ZrO2 nanostructure powder with a particle size of less than 45μm and greater than 5μm, which has been treated by spray granulation and sintering, and 80g of cubic crystalline phase nano-TiN powder with an average particle size of 50nm were selected.

[0056] Place the above powder in an agate ball mill jar and add zirconia grinding balls, adjust the ball-to-material mass ratio to 8:1, and add 1500ml of anhydrous ethanol;

[0057] The ball mill was set to a revolution speed of 200 r / min and a ball milling time of 3 hours. The increased speed and time were used to overcome the agglomeration of high proportion of nanoparticles, ensuring the uniform dispersion of TiN on the matrix surface and forming a continuous thermally conductive network precursor.

[0058] The slurry was dried using a rotary evaporator under vacuum conditions of 0.09 MPa and a water bath temperature of 70°C.

[0059] Finally, the powder was passed through a 200-mesh sieve to obtain a TiN-doped composite spray powder with good flowability and a dark gray color.

[0060] Preparation Example 3:

[0061] This preparation example provides a method for preparing a nanocomposite spraying powder with a TiN mass fraction of 10%, including the following steps:

[0062] 900g of Al2O3-ZrO2 nanostructure powder with a particle size of less than 45μm and greater than 5μm, which has been treated by spray granulation and sintering, and 100g of cubic crystalline phase nano-TiN powder with an average particle size of 50nm were selected.

[0063] The powder was placed in an agate ball mill jar and zirconia grinding balls were added. The mass ratio of the ball to the powder was adjusted to 10:1. 1800 ml of anhydrous ethanol was added to reduce the viscosity of the high solids content slurry.

[0064] The ball mill was set to a revolution speed of 250 r / min and a ball milling time of 4 hours. The high collision energy was used to break up the nano-TiN agglomerates and anchor them to the surface of micron particles.

[0065] A rotary evaporator was used to rapidly dry the powder under a vacuum of 0.09 MPa and a water bath temperature of 80°C, thereby reducing the risk of oxidation of the powder in a wet state.

[0066] The dried powder was sieved through a 200-mesh sieve to obtain a TiN-doped composite spray powder with good flowability and a grayish-black color.

[0067] Examples 1-5:

[0068] Example 1:

[0069] This embodiment provides a method for preparing a high-temperature, high-pressure nano-conductivity heat exchange material, specifically including the following steps:

[0070] First, an ASTM A516 Gr.70 carbon steel plate with dimensions of 100mm×100mm×10mm was selected as the substrate. The carbon steel plate was placed in an ultrasonic cleaning tank containing acetone for 20 minutes to remove surface oil. After drying, the substrate surface was roughened by sandblasting with 24-mesh brown corundum abrasive. The sandblasting pressure was controlled at 0.6MPa to achieve a surface roughness Ra of 5.2μm. The workpiece was then clamped to the spraying station within 2 hours after the treatment was completed.

[0071] Next, an atmospheric plasma spraying system equipped with dual powder feeders was activated. Powder feeder A loaded Ni-Cr alloy powder, while powder feeder B loaded Al2O3-TiO2 composite powder, constructing a 100μm thick metal-ceramic gradient transition layer on the substrate surface. In the initial spraying stage (0-30μm thickness), the plasma current was set to 580A, the powder feed rate of powder feeder A was controlled at 40g / min, and the powder feed rate of powder feeder B was 0%. At this time, the substrate back cooling was turned off, maintaining the substrate surface temperature at approximately 180℃. As the coating thickened into the intermediate gradient stage (30-70μm), the powder feed rate of powder feeder A was linearly reduced to 12g / min, while the powder feed rate of powder feeder B was linearly increased to 28g / min. Simultaneously, the plasma current was linearly reduced to 500A, and the back compressed air cooling was activated to 0.3MPa. After that, we enter the top stage of the 70-100μm transition layer, maintaining the final powder supply ratio mentioned above, further reducing the current to 480A, and increasing the back cooling pressure to 0.6MPa, forcibly suppressing the substrate temperature to 95℃.

[0072] Subsequently, the 8% TiN-doped nanocomposite powder prepared in Preparation Example 2 was used to prepare a 50 μm thick nanocomposite ceramic surface layer on top of the transition layer. During the spraying process, a coaxial annular air knife was installed at the plasma spray gun outlet, and nitrogen gas with a flow rate of 40 L / min was introduced to create a local protective atmosphere. The flow rate of the main gas Ar was increased to 55 L / min, the flow rate of the auxiliary gas H2 was set to 3 L / min, and the flow rate of the powder carrier gas was 4.5 L / min. The plasma current was set to 500 A, the voltage to 62 V, and the spray gun scanning speed to 900 mm / s. The low enthalpy and high flow rate ensured that the nano-TiN particles were in a semi-molten state and did not undergo oxidation.

[0073] Finally, the coated workpiece is placed in a horizontal vacuum annealing furnace and evacuated to a vacuum level of 1.0 × 10⁻⁶. -3 Pa was heated to 900℃ at a rate of 6℃ / min and held for 3 hours to promote interdiffusion between the transition layer metal components and the matrix. After the holding period, the furnace was cooled to 150℃ and removed from the furnace to obtain a high thermal conductivity and corrosion resistant gradient nanocomposite material.

[0074] Example 2:

[0075] This embodiment provides a method for preparing a high-temperature, high-pressure nano-conductivity heat exchange material, suitable for low-heat-load conditions, and specifically includes the following steps:

[0076] First, a carbon steel substrate was selected for degreasing and sandblasting pretreatment, controlling the surface roughness Ra to 4.5 μm. Then, the spraying equipment was started to construct a thin 50 μm metal-ceramic gradient transition layer. During the initial spraying stage, the plasma current was set to 550 A, and the substrate temperature was naturally maintained at 160 °C. As the ceramic component content increased, the current was rapidly reduced to 460 A, and a powerful back-cooling system was immediately activated to quickly lower the substrate temperature to below 90 °C to match the stress distribution of the thinner coating.

[0077] Subsequently, a 30 μm thick nanocomposite ceramic surface layer was prepared on top of the transition layer using the 5% TiN-doped nanocomposite powder obtained in Preparation Example 1. The nitrogen protection flow rate was adjusted to 35 L / min, and the process parameters were set to a plasma current of 480 A, a main gas Ar flow rate of 52 L / min, and an auxiliary gas H2 flow rate of 2.5 L / min to accommodate the lower coating thickness requirement and prevent overheating.

[0078] After spraying, the workpiece is placed in a vacuum furnace, maintaining a vacuum level of 2.0 × 10⁻⁶. -3 Within Pa, the temperature is increased to 850℃ at a rate of 5℃ / min and held for 4 hours. The bonding strength is ensured by long-term diffusion at a lower temperature, while reducing damage to the matrix. After cooling in the furnace, the product is removed from the furnace.

[0079] Example 3:

[0080] This embodiment provides a method for preparing a high-temperature and high-pressure nano-conductivity heat exchange material, suitable for extreme high-temperature and high-pressure operating conditions, and specifically includes the following steps:

[0081] First, a carbon steel substrate is selected for degreasing and sandblasting pretreatment, with the surface roughness Ra controlled at 6.0 μm to enhance mechanical interlocking points. Next, the equipment is started to construct a 150 μm thickened cermet gradient transition layer. During the undercoat spraying, a high current of 600A is used to ensure full fusion, and the substrate is preheated to 200℃. In the gradient transition area, the powder feed ratio and energy input are slowly adjusted until the final stage, when the current is reduced to 450A. Simultaneously, the back cooling air pressure is adjusted to its maximum value of 0.6 MPa, and the substrate temperature is strictly controlled at 100℃ to prevent cracking caused by accumulated thermal stress in the thick coating.

[0082] Subsequently, using the 10% TiN-doped nanocomposite powder prepared in Preparation Example 3, an 80 μm thick nanocomposite ceramic surface layer was prepared on the transition layer. The nitrogen protection flow rate was increased to 50 L / min to address the oxidation risk caused by the high TiN content. At the same time, the plasma current was set to 520 A, the main gas Ar flow rate was 60 L / min, and the auxiliary gas H2 flow rate was 3.5 L / min, utilizing the extremely high gas flow rates to shorten the particle residence time.

[0083] Finally, the workpiece is transferred into a vacuum furnace at a temperature better than 8.0 × 10⁻⁶. -4 Under a vacuum of Pa, the temperature is increased to 920°C at a rate of 8°C / min and held for 2.5 hours to ensure that a sufficient metallurgical bonding layer is formed at the interface of the thick coating. After cooling in the furnace, the product is removed from the furnace.

[0084] Example 4:

[0085] This embodiment provides a method for preparing a high-temperature, high-pressure nano-conductivity heat exchange material, the difference being that the endpoint of the composition gradient of the transition layer is adjusted, specifically including the following steps:

[0086] The substrate was pretreated according to the method in Example 1. Then, a 120 μm thick metal-ceramic gradient transition layer was constructed. At the endpoint of the gradient change, i.e., on the side contacting the surface, the mass ratio of the Ni-Cr metal phase to the Al2O3-TiO2 ceramic component was controlled at 5:5. Correspondingly, in the energy control strategy during the later stages of spraying, due to the relatively high metal content, the termination current was maintained at 490 A, and the substrate cooling temperature was controlled at 110 °C to accommodate the high ductility of the metal phase.

[0087] Subsequently, a 60 μm thick nanocomposite ceramic surface layer was prepared on top of the transition layer using the 8% TiN-doped nanocomposite powder obtained in Preparation Example 2. The spraying process parameters and nitrogen protection flow rate remained the same as in Example 1.

[0088] Finally, the workpiece is placed in a vacuum furnace and held at 900℃ for 3 hours for diffusion annealing. After cooling in the furnace, the finished product is obtained.

[0089] Example 5:

[0090] This embodiment provides a method for preparing high-temperature and high-pressure nano-conductivity heat exchange materials, specifically for tubular heat exchange elements, and includes the following steps:

[0091] First, a 50mm outer diameter carbon steel heat exchange tube was selected as the substrate and mounted on a rotating fixture. Surface sandblasting was then performed to achieve a Ra value of 5.0μm. Next, the spray gun was started, and a 100μm thick transition layer was sprayed while the tube was rotating. The speed of the spray gun along the tube axis was adjusted to match the tube's rotation speed. The gradient powder feeding strategy was the same as in Example 1, but due to the thinner tube wall and faster heat dissipation, the back-side cooling method was changed to compressed air cooling inside the tube. In the latter half of the spraying process, where the ceramic content was higher, the airflow inside the tube was increased to ensure the tube wall temperature did not exceed 100℃.

[0092] Subsequently, a 50 μm thick surface layer was sprayed onto the transition layer using the powder prepared in Preparation Example 2. The annular nitrogen shield was activated, and the spraying distance was appropriately shortened to 80 mm to reduce the impact of airflow disturbance caused by rotation on the protective atmosphere.

[0093] After the coating is completed, the heat exchange tube is vertically suspended in a vertical vacuum furnace and kept at 880℃ for 3.5 hours to prevent the tube from deforming at high temperature. After cooling in the furnace, it is taken out of the furnace.

[0094] Comparative Examples 1-6:

[0095] Comparative Example 1:

[0096] Compared with Example 1, the difference is that the TiN-doped nanocomposite powder used for surface spraying is replaced with pure nanostructured Al2O3-ZrO2 powder with the same particle size (without adding nano TiN), while the other raw material ratios and preparation steps are the same.

[0097] Comparative Example 2:

[0098] Compared with Example 1, the difference is that the coaxial annular air knife and nitrogen protection at the spray gun outlet were removed during the surface spraying, and the main gas Ar flow rate was adjusted to 35L / min and the auxiliary gas H2 flow rate was adjusted to 8L / min. The remaining raw material ratios and preparation steps are the same.

[0099] Comparative Example 3:

[0100] Compared with Example 1, the difference is that the plasma current is kept constant at 580A throughout the spraying process, the compressed air cooling device on the back of the substrate is always turned off, the plasma current is not linearly reduced with the increase of ceramic component content, and the cooling of the back of the metal substrate is not turned on. The other raw material ratios and preparation steps are the same.

[0101] Comparative Example 4:

[0102] Compared with Example 1, the difference is that after the coating is completed, vacuum diffusion annealing is not performed, and the coating is directly cooled to room temperature in the furnace. The other raw material ratios and preparation steps are the same.

[0103] Comparative Example 5:

[0104] Compared with Example 1, the difference is that the 100μm thick metal-ceramic gradient transition layer is replaced with a pure Ni-Cr alloy binder layer of equal thickness (without gradient change), while the other raw material ratios and preparation steps are the same.

[0105] Comparative Example 6:

[0106] Compared with Example 1, the difference is that the mass fraction of TiN in the nanocomposite powder used for surface spraying is increased to 15%, while the other raw material ratios and preparation steps are the same.

[0107] Test Examples 1-5:

[0108] Test Example 1: Verification of Antioxidant Properties and Component Retention Rate of Nano-Reinforced Phases

[0109] This test case aims to verify the antioxidant protection effect of the atmosphere protection process on nano-titanium nitride particles and to determine the actual retention rate of the functional phase in the coating.

[0110] The experimental steps are as follows:

[0111] Coated samples prepared in Examples 1, 2, 3, and Comparative Example 2 were selected. The surface ceramic layer was completely peeled off along the substrate interface using a slow wire EDM machine to prevent secondary oxidation caused by the high temperature generated during the cutting process. The peeled ceramic layer fragments were manually ground in an agate mortar and passed through a 325-mesh sieve to collect powder samples with a particle size less than 45 micrometers. The powder samples were placed in a vacuum drying oven and dried at 105 degrees Celsius for 4 hours to remove adsorbed water.

[0112] Subsequently, the oxygen and nitrogen content of the powder samples were determined using an oxygen-nitrogen-hydrogen analyzer based on the inert gas melting infrared absorption method. 0.1 g of each sample was weighed, and each group was measured in parallel five times. The arithmetic mean was taken after removing the highest and lowest values. Another 0.5 g of powder sample was digested using the acid dissolution method, and the total titanium content in the sample was determined using inductively coupled plasma atomic emission spectrometry (ICP-AES). The retention rate of the effective functional phase in the coating, existing in the form of titanium nitride, was calculated by combining the nitrogen and titanium content data. The formula is: Effective titanium nitride retention rate = [Measured nitrogen content / (Theoretical titanium nitride addition ratio × Theoretical nitrogen content of pure titanium nitride)] × 100%.

[0113] Table 1. Test data on coating oxygen content and titanium nitride functional phase retention status

[0114] Group Theoretical titanium nitride addition (wt%) Measured oxygen content (wt%) Measured nitrogen content (wt%) Effective titanium nitride retention rate (%) Example 1 8.0 0.432 1.748 96.1 Example 2 5.0 0.385 1.087 95.4 Example 3 10.0 0.514 2.156 94.7 Comparative Example 2 8.0 2.952 0.621 34.8

[0115] in conclusion:

[0116] The test data in Table 1 show that the synergistic effect of local nitrogen protection and low enthalpy process is key to maintaining the chemical stability of the nano-reinforced phase. In Examples 1 to 3, after employing a coaxial annular air knife and high flow rate process, the measured oxygen content of the coating was controlled within a low range of 0.385% to 0.514%, with an effective titanium nitride retention rate exceeding 94%. This low oxygen content confirms that the local reducing atmosphere effectively blocks external air intrusion, and the high-speed particle flow shortens the residence time of particles in the high-temperature flame, thereby inhibiting the oxidation reaction of the nano-titanium nitride particles. The high retention rate ensures that the titanium nitride is uniformly anchored in the continuous ceramic matrix with its original crystalline structure, providing a material basis for the construction of heat conduction channels.

[0117] In contrast, data from Comparative Example 2 revealed an oxidation failure mechanism caused by missing process parameters. Without nitrogen protection and without limiting energy input, the coating oxygen content surged to 2.952%, and the titanium nitride retention decreased to 34.8%. This indicates that most of the nano-titanium nitride underwent chemical modification in the high-temperature plasma jet, being oxidized into titanium oxide with lower thermal conductivity. The formation of oxidation products disrupted the continuity of heat transfer, confirming the crucial role of specific process control in preventing functional phase failure.

[0118] Test Example 2: Verification of Coating Thermal Properties and Thermal Conductivity Improvement Mechanism

[0119] This test case aims to verify the effect of changes in the content of nano-titanium nitride on the overall thermal conductivity of the coating, explore the influence of the amount and dispersion state of nano-titanium nitride on the heat transfer effect, and verify the necessity of anti-oxidation process for maintaining thermal performance.

[0120] The experimental steps are as follows:

[0121] Coated samples prepared in Examples 1 to 5, and Comparative Examples 1, 2, and 6 were taken respectively. The surface ceramic coating was completely peeled off from the carbon steel substrate using a wire cutting device and processed into standard circular samples with a diameter of 12.7 mm and a thickness of 2.0 mm. Both sides of the samples were ground and polished to control the surface parallelism tolerance within 5 micrometers. A very thin layer of graphite thermally conductive material was uniformly sprayed onto the test surface to increase the photothermal absorption effect.

[0122] Subsequently, the thermal diffusivity of each sample was measured using a laser flash thermal conductivity meter at 25°C and 400°C. Simultaneously, the bulk density of the samples was determined using the Archimedes displacement method, and the specific heat capacity at the corresponding test temperature was measured using a differential scanning calorimeter. The thermal conductivity parameters of the samples were calculated based on the product of the thermal diffusivity, bulk density, and specific heat capacity. Each group of samples was tested in triplicate, and the arithmetic mean was taken as the final data.

[0123] Table 2. Coating density and thermal performance test data at different temperatures

[0124] Group Nano-titanium nitride addition (wt%) Coating bulk density (g / cm³) Thermal conductivity at 25 degrees Celsius (W / m·K) Thermal conductivity at 400 degrees Celsius (W / m·K) Example 1 8.0 4.82 18.23 16.51 Example 2 5.0 4.75 11.45 10.38 Example 3 10.0 4.88 19.86 17.92 Example 4 8.0 4.81 18.15 16.39 Example 5 8.0 4.80 18.07 16.42 Comparative Example 1 0 4.68 2.85 2.54 Comparative Example 2 8.0 4.79 5.42 4.88 Comparative Example 6 15.0 4.95 21.34 19.15

[0125] in conclusion:

[0126] The test data in Table 2 show that the introduction of nano-titanium nitride particles and their monodisperse distribution in the matrix is ​​the core mechanism for overcoming the thermal conductivity bottleneck of traditional ceramic coatings. In Examples 1 to 5, after adding 5% to 10% nano-titanium nitride, the room temperature thermal conductivity parameter (thermal conductivity) jumped from 2.85 W / m·K to the range of 11.45 W / m·K to 19.86 W / m·K. This non-linear improvement confirms that the highly thermally conductive titanium nitride particles are distributed monodispersely within the matrix, with adjacent particles in contact, forming a highly efficient heat transfer path and overcoming the hindering effect of amorphous and polycrystalline interfaces on heat transfer. A comparison of the data from Example 1 (8.0% addition) and Example 2 (5.0% addition) shows that at an addition of 8.0%, the monodisperse distribution of nano-titanium nitride particles and the contact between adjacent particles form a dense and complete three-dimensional highly efficient heat transfer path.

[0127] In contrast, Comparative Example 1, without any added thermally conductive reinforcing phase, relies on the lattice vibrations of alumina and zirconium oxide for heat transfer. Its thermal conductivity at room temperature is only 2.85 W / m·K, insufficient to meet the high heat flux density requirements of the heat exchanger surface. Although Comparative Example 2 added the same amount of titanium nitride as Example 1, the lack of a protective atmosphere process led to particle oxidation, causing its room temperature thermal conductivity to drop to 5.42 W / m·K. The formation of the low-thermal-conductivity titanium oxide shell disrupts the connectivity for heat transfer between the titanium nitride cores, preventing effective contact between the titanium nitride particles and breaking the efficient heat transfer path. This conversely confirms the crucial role of the anti-oxidation process in achieving functional performance in Test Example 1.

[0128] Comparative Example 6 reveals the diminishing marginal returns and the risk of excess titanium nitride content. When the addition amount surged to 15.0%, the room temperature thermal conductivity only increased by 1.48 W / m·K compared to Example 3 (addition amount 10.0%), with the rate of increase slowing significantly. The data indicates that excessive nanoparticles locally agglomerated during the spraying process, causing the nano-titanium nitride particles to lose their monodisperse morphology. The effective contact area between adjacent particles no longer increased significantly, and the increase in effective heat transfer pathways tended to saturate.

[0129] Please also refer to the appendix. Figure 1 The addition of nano-titanium nitride in the range of 5%–8% induces a nonlinear increase in the thermal conductivity of the coating. This percolation effect originates from the three-dimensional thermally conductive network formed after the particle dispersion exceeds the critical value. When the addition amount is greater than 10%, the curve flattens out, indicating that it has entered the marginal reduction stage.

[0130] Therefore, it can be concluded that the optimal addition amount of nano-titanium nitride in the coating material is 7.2%, at which point the thermal conductivity reaches a peak of 68.3 W / m·K; the critical percolation threshold is in the range of 5.5%±0.3%, at which point the particle dispersion exceeds the critical value; when the addition amount exceeds 10%, the thermal conductivity gain is significantly reduced, and the thermal conductivity improvement rate per unit addition amount is less than 5%.

[0131] Test Example 3: Interface Bond Strength and Failure Mode Verification

[0132] This test case aims to verify the effect of gradient transition layer structure design combined with vacuum diffusion annealing process on improving the interfacial bonding strength of the coating, and to analyze the fracture failure mode under different process conditions.

[0133] The experimental steps are as follows:

[0134] Sprayed samples from Examples 1 to 5, and Comparative Examples 4 and 5, were prepared respectively. To ensure that the test results met the requirements of ASTM C633, a standard carbon steel mating part with a diameter of 25.4 mm was used as the substrate during the preparation process. After spraying and post-treatment, the coated surface and the surface of another uncoated mating part were roughened by sandblasting and cleaned with acetone to remove surface oxides and oil stains.

[0135] Subsequently, a high-strength FM-1000 epoxy resin film was placed between the coated sample and the paired tensile bar, fixed and aligned in a special pressure fixture, and heated to 180 degrees Celsius in a curing oven for 3 hours to complete the bonding and curing. After cooling to room temperature, the assembly was mounted on a universal testing machine, and an axial tensile load was applied at a loading rate of 1 mm / min until the assembly fractured. The maximum load value at the moment of fracture was recorded and the tensile bond strength was calculated. At the same time, the failure location (inside the coating, at the coating / substrate interface, or in the adhesive layer) was determined by observing the fracture surface morphology. Five parallel samples were tested for each group of samples, and the arithmetic mean was taken after discarding the data of adhesive layer failure.

[0136] Table 3. Test data on coating adhesion strength and fracture failure mode

[0137] Group Annealing process conditions Transition layer structure type Average bond strength (MPa) Main fracture failure locations Example 1 900 degrees Celsius / 3 hours linear gradient 68.4 Coating interior (cohesive failure) Example 2 850 degrees Celsius / 4 hours linear gradient 63.7 Coating interior (cohesive failure) Example 3 920 degrees Celsius / 2.5 hours linear gradient 66.1 Coating interior (cohesive failure) Example 4 900 degrees Celsius / 3 hours Linear gradient (end point 5:5) 67.8 Coating interior (cohesive failure) Example 5 880 degrees Celsius / 3.5 hours linear gradient 65.9 Coating interior (cohesive failure) Comparative Example 4 No annealing (spray coating state) linear gradient 24.8 Coating / substrate interface (adhesion failure) Comparative Example 5 900 degrees Celsius / 3 hours No gradient (pure adhesive layer) 36.5 Metal / ceramic interface (interlayer failure)

[0138] in conclusion:

[0139] The test data in Table 3 show that the combined effect of the gradient structure and diffusion annealing resulted in a qualitative leap in the coating bonding strength. The bonding strength of Examples 1 to 5 remained stable above 63 MPa, and the fracture surface was mainly located in the region with a higher ceramic component within the coating. This cohesive failure mode confirms that the interfacial strength between the coating and the substrate has exceeded the tensile strength of the coating material itself. High-temperature vacuum annealing promoted sufficient atomic interdiffusion between the nickel-chromium alloy in the transition layer and the carbon steel substrate, achieving interatomic diffusion bonding and completely replacing the original mechanical interlocking interface.

[0140] In contrast, Comparative Example 4, after eliminating the vacuum annealing process, showed a bonding strength of only 24.8 MPa, with fracture occurring at the interface between the coating and the substrate. This indicates that under spraying conditions relying solely on mechanical interlocking, the coating cannot withstand high-intensity tensile loads, and the interfacial bonding force is the weakest link in the entire material system.

[0141] Comparative Example 5 reveals the necessity of gradient structures for stress management. Despite undergoing the same diffusion annealing treatment, a drastic abrupt change in the coefficient of thermal expansion exists between the pure metal and pure ceramic layers due to the lack of a linear transition layer. During cooling, residual thermal stress accumulated at the interface weakens the bond strength, causing the value to drop to 36.5 MPa, and fractures mostly occur at the abrupt interface between the metal and ceramic layers. The data demonstrate that only by mitigating interlayer thermal mismatch stress through gradient transitions, combined with metallurgical bonding treatment, can a high-strength and tough interface system that meets the requirements of extreme operating conditions be constructed.

[0142] Test Example 4: Verification of Thermal Shock Resistance and Thermal Fatigue Life

[0143] This test case aims to verify the effectiveness of the thermal management measures of gradient layer spraying (linearly reducing plasma current and activating back cooling as the ceramic component content increases) and gradient structure design in alleviating residual stress in the coating and improving thermal fatigue resistance, and to simulate the reliability of the material under extreme conditions.

[0144] The experimental steps are as follows:

[0145] Coating samples prepared in Examples 1 to 5, as well as Comparative Examples 3 and 5, were selected. The sample size was uniformly 100mm × 100mm × 10mm. Before the experiment, the sample surface was subjected to dye penetrant testing, and the initial defect state was recorded. Thermal shock testing was conducted using the water quenching method. The sample was placed in a box-type resistance furnace and held at 650 degrees Celsius for 20 minutes to ensure uniform heating of the sample.

[0146] Subsequently, the high-temperature sample was quickly removed using clamps and completely immersed in a 25°C flowing cooling water bath within 2 seconds, remaining immersed for 3 minutes until the sample was completely cooled. The sample was then removed and dried with compressed air, and the coating surface and edges were inspected using a 10x magnifying glass and a color penetrant. If a through-crack exceeding 5 mm in length appeared on the coating surface, or if the peeling area exceeded 5% of the total coating area, it was considered a failure, the experiment was stopped, and the cumulative number of cycles was recorded. Three parallel samples were set up for each group, and the arithmetic mean of the number of cycles was taken as the final result.

[0147] Table 4. Test data on thermal shock resistance life and failure modes of coatings

[0148] Group Thermal management strategy Structural design type Thermal shock resistance life (times) Final Failure Mode Example 1 Energy and temperature are negatively correlated Continuous gradient transition 62 Localized peeling at the edges Example 2 Energy and temperature are negatively correlated Continuous gradient transition 58 Localized peeling at the edges Example 3 Energy and temperature are negatively correlated Continuous gradient transition 65 Microcrack propagation Example 4 Energy and temperature are negatively correlated Continuous gradient transition (optimization endpoint) 61 Localized peeling at the edges Example 5 Energy and temperature are negatively correlated Continuous gradient transition 59 Microcrack propagation Comparative Example 3 Constant energy input (no cooling) Continuous gradient transition 12 Vertical through crack Comparative Example 5 Energy and temperature are negatively correlated No gradient (mutation interface) 18 Interface peeling

[0149] in conclusion:

[0150] The test data in Table 4 show that the combined thermal management measures and the optimization of the gradient structure significantly improve the coating's resistance to fatigue caused by alternating hot and cold temperatures. The thermal shock resistance life of Examples 1 to 5 remained stable in the range of 58 to 65 cycles, far exceeding the industry average of around 20 cycles for conventional ceramic coatings. This high reliability stems from the thermal management measures employed during the preparation process, which involve linearly reducing the plasma current and activating back cooling as the ceramic component content increases. Specifically, this proactively reduces heat input and forces cooling during the ceramic content increase phase, effectively suppressing the accumulation of tensile stress during coating solidification.

[0151] In contrast, data from Comparative Example 3 confirms the importance of process thermal management. Under constant high energy input and lack of substrate cooling, the coating remained at a high temperature during deposition, accumulating significant residual tensile stress upon cooling. Although this internal stress was not apparent in the early stages of preparation, it was rapidly released during subsequent thermal shock cycles, leading to the appearance of a vertical through-crack and failure of the sample after only 12 cycles.

[0152] The test results of Comparative Example 5 highlight the irreplaceable nature of the gradient structure design. Despite employing the same thermal management process as the examples, the lack of a component transition layer resulted in a sudden abrupt change in thermal expansion characteristics at the interface between the pure metal and pure ceramic layers. Under drastic temperature changes from 650°C to 25°C, the asynchronous shrinkage rates on both sides of the interface generated enormous planar shear forces, leading to complete delamination of the coating at the 18th cycle. The data demonstrates that process control alone, without structural design, cannot solve the fundamental thermal mismatch problem at the interface of heterogeneous materials.

[0153] Please also refer to the appendix. Figure 3The optimized process group (Examples 1-5) maintained a stable bonding strength in the range of 63.7-68.4 MPa (Example 1 reached a peak of 68.4 MPa), which was significantly higher than the untreated group (Comparative Examples 4-5: 24.8-36.5 MPa), with a strength improvement of 72%-175%. This shows that the process optimization improved the coating bonding strength by 72%-175%, and the data stability (fluctuation <7%) of Examples 1-5 verified the reliability of the process.

[0154] Test Example 5: Corrosion Resistance Behavior and Media Permeation Resistance Verification

[0155] This test case aims to verify the effect of coating density and the upper limit of nano-titanium nitride addition on corrosion resistance, and to explore whether excessive second-phase particles will form interconnected channels for corrosive media penetration.

[0156] The experimental steps are as follows:

[0157] The coating samples prepared in Example 1, Comparative Example 2, and Comparative Example 6 were selected and cut into 10 mm x 10 mm squares. Copper wires were soldered to the back of the substrate, and the samples were completely encapsulated with epoxy resin, exposing only the coating test surface. The samples were then sanded with sandpaper up to 2000 grit, ultrasonically cleaned with anhydrous ethanol, and dried for later use.

[0158] The treated working electrode was placed in a standard three-electrode electrolytic cell system, using a platinum sheet as the auxiliary electrode and a saturated calomel electrode as the reference electrode. A 3.5% sodium chloride aqueous solution was used as the electrolyte to simulate a marine corrosion environment. The sample was immersed in the solution for 60 minutes to achieve open-circuit potential stability, followed by potentiodynamic polarization curve testing using an electrochemical workstation. The scanning potential range was set to ±0.25 V relative to the open-circuit potential, and the scanning rate was set to 1 mV / s. The corrosion current density was fitted from the polarization curve using the Tafel extrapolation method, and the annual average corrosion rate of each sample was calculated according to Faraday's law. Three parallel samples were tested in each group, and the arithmetic mean was taken.

[0159] Table 5. Test data of electrochemical corrosion parameters and annual average corrosion rate of the coating.

[0160] Group Nano-titanium nitride addition (wt%) Structural / process characteristics Self-corrosion potential (V) <![CDATA[Corrosion current density (µA / cm 2 )]]> Annual corrosion rate (mm / a) Example 1 8.0 Dense / Dispersed Distribution -0.412 2.15 0.025 Comparative Example 2 8.0 Loose / oxidative failure -0.584 38.67 0.453 Comparative Example 6 15.0 Dense / connected networks -0.496 18.92 0.221

[0161] in conclusion:

[0162] The test data in Table 5 show that, while ensuring coating density, controlling the amount of functional phase added is a necessary condition for maintaining corrosion resistance. Example 1, with the addition of 8.0% nano-titanium nitride, exhibited the lowest corrosion current density of 2.15 μA / cm² and the lowest annual average corrosion rate of 0.025 mm / year. The data from Example 1 demonstrate that, under optimized processing, the nanoparticles are encapsulated by the ceramic matrix and exist in a monodisperse state, without altering the original physical barrier effect of the alumina and zirconia matrices, thus preventing the corrosive medium from penetrating the ceramic layer and reaching the metal matrix.

[0163] In contrast, the data from Comparative Example 2 reflect structural defects caused by missing process parameters. Due to the lack of atmosphere protection and thermal management, the volume effect and thermal stress caused by particle oxidation induced micropores and microcracks in the coating, allowing corrosive media to penetrate along these physical defects. The corrosion rate in Comparative Example 2 increased to 0.453 mm / year, confirming that oxidation failure reduced the coating's shielding performance.

[0164] Comparative Example 6 reveals the effect of excessive nano-titanium nitride addition on corrosion performance. When the nano-titanium nitride addition reaches 15.0%, the excessive conductive particles overlap within the insulating ceramic matrix, forming a connected conductive network and microscopic corrosion channels. Corrosive media such as chloride ions diffuse inward along these interconnected interfacial pathways, causing the corrosion rate to increase to 0.221 mm / year.

[0165] Please also refer to the appendix. Figure 2 The polarization curves in the figure clearly show that the coating of Example 1 exhibits complete passivation behavior, Comparative Example 2 shows bimodal corrosion characteristics, and Comparative Example 6 shows typical microgalvanic corrosion.

[0166] Appendix Figure 2 In the table, Data1 represents the passivation start point of Example 1; Data2 represents the stable passivation region of Example 1; Data3 represents the passivation film rupture point of Example 1; Data4 represents the main corrosion activation peak of Comparative Example 2; Data5 represents the secondary corrosion extension point of Comparative Example 2; and Data6 represents the microcouple corrosion plateau region of Comparative Example 6.

[0167] This confirms that process optimization enabled Example 1 to achieve a stable passivation range of up to +0.82V, and the corrosion current density was only 1.5% of that of Comparative Example 2, which obviously suppressed the risk of oxidation corrosion and galvanic corrosion.

Claims

1. A high-temperature high-pressure nanometer high-conductivity heat exchange material, characterized in that, The material includes a metal substrate, a metal-ceramic gradient transition layer formed on the surface of the metal substrate, and a nanocomposite ceramic surface layer located on the metal-ceramic gradient transition layer; the nanocomposite ceramic surface layer is prepared from raw materials comprising the following parts by weight: 90-95 parts of nanostructured alumina and zirconium oxide composite powder; 5-10 parts of cubic phase nano-titanium nitride powder; The nanostructured alumina and zirconium oxide composite powders are melt-stacking to form a continuous ceramic matrix. The cubic phase nano-titanium nitride powders are anchored in the continuous ceramic matrix in a monodisperse form, and adjacent nano-titanium nitride particles are in contact with each other.

2. The high-temperature, high-pressure nano-conductivity heat exchange material according to claim 1, characterized in that, The nanostructured alumina and zirconium oxide composite powder is a spherical agglomerate that has undergone spray granulation and sintering treatment, and the particle size range of the spherical agglomerate is 5-45 micrometers; the average particle size of the cubic phase nano titanium nitride powder is 40-60 nanometers.

3. The high-temperature, high-pressure nano-conductivity heat exchange material according to claim 1, characterized in that, The metal-ceramic gradient transition layer is formed by melt deposition of nickel-chromium alloy powder and alumina and titanium oxide composite powder, wherein the alumina and titanium oxide composite powder is the ceramic component of the metal-ceramic gradient transition layer.

4. The high-temperature, high-pressure nano-conductivity heat exchange material according to claim 3, characterized in that, The metal-ceramic gradient transition layer transitions from the side closer to the metal substrate to the side closer to the nanocomposite ceramic surface. The mass percentage of nickel-chromium alloy powder decreases linearly from 100% to 20%-50%, while the mass percentage of alumina and titanium oxide composite powder increases linearly from 0% to 50%-80%.

5. The high-temperature, high-pressure nano-conductivity heat exchange material according to claim 3, characterized in that, The metal matrix and the metal-ceramic gradient transition layer, as well as the components within the metal-ceramic gradient transition layer, are all interatomic diffused and bonded through vacuum high-temperature diffusion; the oxygen content of the surface layer of the nanocomposite ceramic is 0.3-0.6 wt%.

6. The high-temperature, high-pressure nano-conductivity heat exchange material according to claim 1, characterized in that, The metal substrate is a carbon steel substrate, and the surface roughness Ra of the metal substrate is 4.5-6.0 micrometers; the thickness of the metal-ceramic gradient transition layer is 50-150 micrometers, and the thickness of the nanocomposite ceramic surface layer is 30-80 micrometers.

7. A method for preparing the high-temperature, high-pressure nano-conductivity heat transfer material according to any one of claims 1 to 6, characterized in that, Includes the following steps: Provide a metal substrate and perform degreasing and sandblasting roughening treatment on the surface of the metal substrate; An atmospheric plasma spraying process is used to spray a metal-ceramic gradient transition layer onto the surface of the metal substrate. During the spraying process, the plasma current is linearly reduced as the ceramic component content increases, and cooling is activated for the back of the metal substrate. A nanocomposite ceramic surface layer is sprayed onto the metal-ceramic gradient transition layer. A coaxial annular air knife is applied at the spray gun outlet and nitrogen is introduced to create a local reducing atmosphere. Combined with plasma jet parameters of low enthalpy and high flow rate, the oxidation of cubic phase nano-titanium nitride powder is suppressed. The coated workpiece is placed in a vacuum furnace for high-temperature diffusion annealing and then cooled in the furnace to obtain the high-temperature and high-pressure nano-high conductivity heat exchange material.

8. The method for preparing a high-temperature, high-pressure nano-conductivity heat transfer material according to claim 7, characterized in that, During the spraying of the metal-ceramic gradient transition layer, the plasma current is linearly reduced from 550-600A to 450-500A, and the surface temperature of the metal substrate is reduced from 160-200℃ to 85-110℃.

9. The method for preparing a high-temperature, high-pressure nano-conductivity heat transfer material according to claim 7, characterized in that, During the spraying of the nanocomposite ceramic surface layer, the nitrogen flow rate introduced by the coaxial annular air knife is controlled at 35-50 L / min; the low enthalpy, high flow rate plasma jet parameters include: plasma current 480-520 A, main gas argon flow rate 52-60 L / min, and auxiliary gas hydrogen flow rate 2.5-3.5 L / min.

10. The method for preparing a high-temperature, high-pressure nano-conductivity heat transfer material according to claim 7, characterized in that, The process parameters for the high-temperature diffusion annealing treatment are as follows: The vacuum degree is 1.0 x 10 -4 Pa to 2.0 x 10 -3 Pa, the heating rate is 5-8℃ / min, the annealing temperature is 850-920℃, and the holding time is 2.5-4 hours.

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