Modified carbon fiber material for unmanned aerial vehicle and processing technology of modified carbon fiber material

Abstract

The invention relates to the field of unmanned aerial vehicle materials, and discloses a modified carbon fiber material for an unmanned aerial vehicle, which comprises the following structural layers: a carbon fiber matrix; the Ni-Ti transition layer is deposited on the surface of the base body, the thickness of the Ni-Ti transition layer is 18-22 microns, and the mass ratio of Ni to Ti is 49.5: 50.5-50.5: 49.5; the nano composite coating covers the transition layer and comprises a nickel aminosulfonate matrix, nano alpha-Al2O3 and silicon carbide whiskers, the nano alpha-Al2O3 and the silicon carbide whiskers are dispersed in the nickel aminosulfonate matrix, the content of Al2O3 is 78-82 g / L, and the content of the silicon carbide whiskers is 14-16 g / L. Through the gradient composite interface structure design, cross-scale synergistic reinforcement of the metal matrix, the ceramic layer and the resin layer is achieved, the synergistic effect of chemical bonding and physical anchoring is further enhanced, and the material can still keep excellent interface stability in a thermal-mechanical coupling environment.

Description

Technical Field

[0001] The present invention relates to the technical field of unmanned aerial vehicle (UAV) materials, and in particular to a modified carbon fiber material for UAVs and a processing technology thereof. Background Art

[0002] Carbon fiber composites, due to their high specific strength and low thermal expansion coefficient, have become a core material for lightweight UAV structural components. Existing technologies focus on carbon fiber surface modification, the introduction of composite reinforcements, and interface optimization, involving processes such as plasma activation, magnetron sputtering coating, micro-arc oxidation ceramicization, and resin-matrix hybridization. Related research focuses on enhancing a single property, such as improving wear resistance through nanoparticle dispersion strengthening or mitigating interfacial thermal stress using gradient coating designs.

[0003] The current technology system has multiple bottlenecks in the synergistic reinforcement of carbon fiber-metal-resin cross-interface: existing carbon fiber composite materials in UAV applications still face problems such as interlayer cracking caused by insufficient interface bonding strength, thermal stress failure caused by differences in thermal expansion coefficients of heterogeneous materials, decreased erosion resistance caused by uneven dispersion of nano-reinforced phases, and poor batch stability caused by mismatch of multi-process process parameters, which seriously restrict their reliability and service life in extreme environments. Summary of the Invention

[0004] The purpose of the present invention is to provide a modified carbon fiber material for drones and its processing technology, which solves the problems in the prior art such as interlayer cracking caused by insufficient interface bonding strength, thermal stress failure caused by mismatch in thermal expansion coefficients of heterogeneous materials, insufficient erosion resistance caused by uneven dispersion of nano-reinforced phases, and poor batch stability caused by poor synergy of multi-process parameters.

[0005] To achieve the above objectives, the present invention is implemented through the following technical solutions: A modified carbon fiber material for a drone, comprising the following structural layers: carbon fiber matrix; High-modulus carbon fiber braids or prepregs are used as the primary load-bearing structure. The substrate surface undergoes plasma activation pretreatment to create a micro-nano roughened structure and introduce oxygen-containing functional groups. This treatment does not alter the mechanical properties of the carbon fiber itself, but provides active sites for chemical bonding of the subsequent metal layer, avoiding the fiber damage associated with traditional mechanical polishing.

[0006] The Ni-Ti transition layer deposited on the substrate surface has a thickness of 18 to 22 μm and a mass ratio of Ni to Ti of 49.5:50.5 to 50.5:49.5; The quasi-equiatomic ratio of Ni-Ti alloy makes it have both high ductility and moderate thermal expansion coefficient, which effectively alleviates the thermal expansion between carbon fiber (CTE≈0) and subsequent metal coating (Ni-based coating, CTE≈13×10 -6 / ℃) between the thermal mismatch stress.

[0007] The amorphous phase (accounting for 60-70%) can inhibit crack propagation, and the nanocrystalline phase (size <50nm) provides load-bearing capacity. The composite structure of the two enables the transition layer to maintain interface integrity under thermal cycling loads.

[0008] The nano-composite coating covering the transition layer comprises a nickel sulfamate matrix and nano-α-Al2O3 and silicon carbide whiskers dispersed therein, wherein the Al2O3 content is 78-82 g / L and the silicon carbide whisker content is 14-16 g / L; Nano-α-Al2O3 (hardness HV ≥ 2000) acts as a hard reinforcing phase to improve the wear resistance of the coating; silicon carbide whiskers (aspect ratio > 20) toughen the coating through bridging and pull-out mechanisms, inhibiting brittle cracking of the coating.

[0009] A nickel sulfamate system (instead of Watt nickel) is used to reduce the internal stress of the plating solution. Combined with the current reversing effect of the three-pulse power supply, the nanoparticles are forced to be oriented in the coating, avoiding the agglomeration defects of conventional DC electroplating.

[0010] The epoxy-silane hybrid protective layer coated on the surface of the composite coating comprises 94-96 wt% of epoxy resin E51, 7.5-8.5 wt% of APTES and 4.8-5.2 wt% of POSS.

[0011] The amino groups in APTES react with the epoxy groups, while its siloxane hydrolysis products condense with the metal hydroxyl groups on the surface of the coating to achieve chemical bonding at the resin-metal interface.

[0012] The POSS nanocage structure (particle size 20-30nm) is embedded in the epoxy network as a physical cross-linking point, inhibiting crack propagation through the steric hindrance effect, while improving the heat resistance of the resin layer (Tg increased by ≥30℃).

[0013] Preferably, the Ni-Ti transition layer is prepared by magnetron sputtering with a sputtering power of 280 to 320 W, a substrate temperature of 180 to 220° C., and a deposition rate of 0.45 to 0.55 nm / s.

[0014] The sputtering power was set between 280 and 320 W, a range determined by experiments balancing the target ionization rate with the film quality. When the power was below 280 W, insufficient plasma density resulted in a low deposition rate (<0.4 nm / s), and incompletely ionized metal clusters were likely to form in the film, leading to microscopic pore defects. When the power was above 320 W, although the deposition rate increased, the thermal stress generated by high-energy particles bombarding the substrate could induce microcracks in the film, and the excessive sputtering kinetic energy could damage the plasma-activated layer on the carbon fiber surface.

[0015] The power window of 280 to 320 W ensures effective ionization of the Ni-Ti target, giving the deposited particles moderate kinetic energy (10 to 20 eV), ensuring film density while avoiding structural damage to the carbon fiber matrix. Within this power range, the sputtering yield ratio of Ni to Ti atoms remains stable at 0.98 to 1.02, accurately maintaining the equiatomic ratio of the alloy components.

[0016] In the temperature range of 180-220℃, the deposited atoms have a moderate surface diffusion ability, which promotes the coordinated growth of amorphous phase and nanocrystalline phase (grain size <50nm). The amorphous phase (accounting for 60-70%) absorbs the interfacial stress through the disordered structure, while the nanocrystalline phase provides the load-bearing capacity. The composite film structure formed by the two can effectively inhibit the initiation of cracks. At the same time, the difference in thermal expansion coefficient with the carbon fiber matrix (ΔCTE <4×10 -6 / ℃) is controlled within the safety threshold.

[0017] Preferably, the modified carbon fiber material for drones is characterized in that, in the nano-composite coating, the particle size of nano α-Al3O3 is 45-55 nm, the diameter of silicon carbide whiskers is 180-220 nm, and the coating thickness is 78-82 μm.

[0018] The particle size of nano-α-Al2O3 is strictly limited to the range of 45 to 55 nm. This size is determined by calculating the equilibrium relationship between the surface energy of the nanoparticles and their Brownian motion ability in the plating solution. When the particle size is less than 45 nm, the surface energy of the particles increases significantly, causing them to irreversibly agglomerate in the nickel sulfamate plating solution due to van der Waals forces, forming micron-sized particle clusters. When the particle size is greater than 55 nm, the specific surface area of the particles decreases, reducing the contact interface with the nickel substrate and weakening the dispersion strengthening effect.

[0019] During the pulse electroplating process, the surface charge (Zeta potential ≈ +25mV) of α-Al2O3 particles with a particle size range of 45 to 55nm reacts with the aminosulfonate ions (-SO3 -) forms an electrostatic repulsion effect, effectively inhibiting particle agglomeration. At the same time, the sedimentation rate of particles of this size (calculated by Stokes' formula is 0.02-0.03 mm / s) matches the plating solution circulation flow rate (0.5-0.8 m / s), ensuring their uniform suspension in the plating solution.

[0020] The diameter of the silicon carbide whiskers is set to 180-220nm, which is based on the principle of matching the bending strength of the whiskers with the internal stress of the coating. When the diameter is less than 180nm, the whiskers are easily broken by the shear force of the fluid during the electroplating process and lose their toughening effect. When the diameter is greater than 220nm, the difference in thermal expansion coefficient between the whiskers and the nickel substrate (ΔCTE≈4×10 -6 / ℃) leads to interface stress concentration, which becomes the source of crack initiation.

[0021] The coating thickness is controlled between 78 and 82 μm, a range determined by the stress gradient distribution within the coating and functional requirements. When the thickness is less than 78 μm, the coating's coverage of the substrate is insufficient, and localized corrosion occurs due to interface defects. When the thickness is greater than 82 μm, the residual stress within the coating (primarily due to the thermal expansion difference between the nickel substrate and the reinforcement phase) exceeds a critical value (>300 MPa), causing the coating to warp or peel.

[0022] Preferably, the curing process of the epoxy-silane hybrid protective layer includes three stages: Stage 1: Curing at 78-82°C for 1.9-2.1h, vacuum degree -0.094-0.096MPa; Curing takes place at 78-82°C and a vacuum of -0.094 to -0.096 MPa for 1.9 to 2.1 hours. This stage is dominated by initial crosslinking of the resin and interfacial wetting. The vacuum environment forces the solvent and low-molecular-weight volatiles (residual monomers and water molecules) in the resin to transform from liquid to gaseous form, where they are removed by negative pressure. The temperature is controlled above the epoxy resin's glass transition temperature (Tg≈65°C) but below its rapid curing onset temperature (≈90°C) to ensure adequate mobility of the molecular segments.

[0023] The second stage: curing at 118-122°C for 0.95-1.05h, vacuum degree -0.094-0.096MPa; The temperature is raised to 118-122°C, while maintaining a vacuum of -0.094 to -0.096 MPa, and the curing time is 0.95 to 1.05 hours. During this phase, the epoxy resin enters a rapid crosslinking phase, accelerating the ring-opening reaction between APTS and epoxy groups. Simultaneously, the POSS nanoparticles begin to interpenetrate with the resin network. The vacuum environment continuously suppresses bubble regeneration and promotes the removal of crosslinking byproducts (water and alcohols).

[0024] The third stage: curing at 148-152°C for 0.48-0.52h, normal pressure environment.

[0025] Curing at 148-152°C under normal pressure for 0.48-0.52 hours completes deep crosslinking of the resin network and relaxation of residual stresses. Ambient pressure allows oxygen to participate in oxidative crosslinking of the resin surface, forming a dense surface layer. Temperatures exceeding the ultimate Tg of the resin (≈140°C) allow for full molecular chain movement to eliminate internal defects.

[0026] Preferably, a porous ceramic layer generated by micro-arc oxidation is provided between the nano-composite coating and the Ni-Ti transition layer, with a porosity of 30-40% and a thickness of 8-12 μm.

[0027] The porosity of the porous ceramic layer is limited to 30-40%, a range based on the balance between pore structure and interfacial bonding strength and stress distribution. When the porosity is below 30%, the ceramic layer exhibits a dense structure, and its elastic modulus (≈200 GPa) differs significantly from that of the Ni-Ti transition layer (≈80 GPa), leading to interfacial stress concentration. When the porosity is above 40%, although the elastic modulus decreases (≈120 GPa), the excessive pore connectivity weakens the mechanical anchoring effect between the coating and the transition layer.

[0028] The thickness of the ceramic layer is set to 8 to 12 μm, which is determined by the comprehensive requirements of interface bonding strength and electrical conductivity. When the thickness is less than 8 μm, the stress buffering capacity of the porous structure is insufficient, and the high voltage discharge (>400V) during the micro-arc oxidation process is easy to break through the transition layer; when the thickness is greater than 12 μm, the insulation of the ceramic layer (resistivity>10 6 Ω·cm) results in the need to apply excessively high voltage (>15V) during subsequent electroplating, causing the plating solution to decompose.

[0029] At a thickness of 8 to 12 μm, the ceramic layer exhibits a gradient structure: Near the transition layer (0-3 μm): The porosity is low (≈35%), rich in Al2O3-SiO2 amorphous phase, and chemically bonded with the Ni-Ti layer through a diffusion interface (width ≈0.5 μm); Middle region (3-8 μm): The porosity gradient rises to 40%, forming a dendritic pore network, which provides a path for the electroplating solution to penetrate; Surface area (8-12 μm): The pore opening diameter is expanded to 2-5 μm, and the surface is covered with nano-scale oxide particles (particle size ≈ 50 nm), increasing the contact area with the nano-composite coating.

[0030] The porous ceramic layer is produced through the synergistic combination of a specific electrolyte (9.5-10.5 g / L Na₂SiO₃ + 4.8-5.2 g / L KOH) and electrical parameters (voltage 440-460 V, frequency 480-520 Hz). A high-voltage pulsed electric field induces plasma discharge on the surface of the Ni-Ti layer, leading to a melting-quenching cycle to form the porous oxide.

[0031] The pulse frequency of 480-520 Hz controls the duration of a single discharge to 10-20 μs and limits the molten pool size to 5-8 μm, thus avoiding pore closure caused by excessive sintering. SiO3 in electrolyte 2- At high discharge temperatures (>3000K), it decomposes into SiO2, which forms an amorphous-nanocrystalline composite ceramic phase with Al2O3 and TiO2 generated by oxidation of the Ni-Ti layer. During the micro-arc oxidation process, the Ni element on the surface of the Ni-Ti layer is selectively oxidized to form NiO (content ≈ 15%), which forms a coherent interface with the Ni matrix of the subsequent nano-composite coating, reducing the interfacial energy barrier.

[0032] A processing technology based on the modified carbon fiber material comprises the following steps: (a) Plasma activation treatment of carbon fiber matrix; The carbon fiber surface was bombarded with Ar / O2 mixed gas (volume ratio 4:1±0.2) at a radio frequency power of 450-550W for 14-16 min, and the vacuum was maintained at 8×10 -3 ~1.2×10 -2 Pa. Nanoscale grooves (Ra = 0.8-1.2 μm) are formed on the surface of the treated fiber and oxygen-containing functional groups (-COOH, -OH) are introduced.

[0033] High-energy Ar + Ions selectively etch along the (002) direction of the carbon fiber graphite crystal plane, forming nano-grooves with an angle of 45° to 60° with the fiber axis, increasing the specific surface area (by about 5 times); The active oxygen free radicals in the O2 plasma react with the carbon surface to generate carboxyl and hydroxyl groups, with a concentration of 8 to 12 per nm. 2 , providing anchor points for subsequent chemical bonding of metal layers.

[0034] (b) Deposition of a Ni-Ti transition layer on the activated surface by magnetron sputtering; A Ni50Ti50 alloy layer is sputtered on the surface of the activated carbon fiber with a power of 280-320 W, a substrate temperature of 180-220° C., a deposition rate of 0.45-0.55 nm / s, and a final thickness of 18-22 μm.

[0035] The substrate temperature of 180-220℃ limits the mobility of sputtered particles on the surface, forming a composite film with an amorphous matrix (accounting for 60-70%) encapsulating nanocrystalline (size <50nm). Its thermal expansion coefficient (CTE = 10-12×10 -6 / ℃) between carbon fiber (CTE≈0) and subsequent coating (CTE≈14×10 -6 / ℃) to achieve thermal stress gradient transition; The Ni / Ti sputtering yield ratio is stabilized at 0.98-1.02 by coordinated regulation of power and gas pressure, ensuring that the alloy composition deviation is <±0.5%.

[0036] (c) performing micro-arc oxidation on the surface of the transition layer to form a porous ceramic layer; A pulse voltage of 440-460V is applied to the surface of the Ni-Ti layer, and the electrolyte is Na2SiO3 (9.5-10.5g / L) + KOH (4.8-5.2g / L). The treatment is carried out for 9-11 minutes to generate an Al2O3-SiO2 composite ceramic layer with a thickness of 8-12μm and a porosity of 30-40%.

[0037] The pulse frequency of 480-520 Hz makes the single discharge duration ≤ 20 μs, the molten pool size is limited to 5-8 μm, and a composite structure of primary pores with a diameter of 1-3 μm and secondary pores with a diameter of 50-200 nm is formed; SiO3 in electrolyte 2- It decomposes into SiO2 at high discharge temperature (>3000K), and forms a ceramic layer of nanocrystalline phase (size 20-50nm) wrapped by an amorphous phase matrix (accounting for 70-80%) with Al2O3 and TiO2 generated by oxidation of Ni-Ti layer. Its surface hydroxyl density reaches 15-20 / nm. 2 .

[0038] (d) preparing a nanocomposite coating on the porous ceramic layer using a three-pulse electroplating process; Pulse electroplating is performed on the surface of the micro-arc oxide layer. The plating solution contains nickel sulfamate (375-385 g / L), nano α-Al2O3 (78-82 g / L), and silicon carbide whiskers (14-16 g / L). A three-pulse mode with a forward pulse width of 0.95-1.05 ms, a reverse pulse width of 0.18-0.22 ms, and an interval of 0.48-0.52 ms is adopted, and the current density J=4.0-0.05C_{Al2O3} is dynamically adjusted.

[0039] High current density of forward pulse (4A / dm 2 ) drives the nanoparticles to adsorb to the cathode surface, and the reverse pulse (0.2ms) dissolves the weakly adsorbed particles, which increases the bonding strength of the particles embedded in the coating by 2 to 3 times; The coupling effect of cathode rotation (28-32 rpm) and ultrasonic waves (36-40 kHz) breaks the boundary layer of the plating solution, making the distribution of nanoparticles in the pores of the micro-arc oxidation layer uniform (CV value <5%).

[0040] (e) Spraying epoxy-silane hybrid resin on the surface of the composite coating and vacuum curing.

[0041] A hybrid resin containing epoxy resin E51 (94-96wt%), APTES (7.5-8.5wt%), and POSS (4.8-5.2wt%) is sprayed onto the surface of the coating with a film thickness of 0.48-0.52mm, and a protective layer is formed through three-stage vacuum curing (78-82℃ / 1.9-2.1h→118-122℃ / 0.95-1.05h→148-152℃ / 0.48-0.52h).

[0042] The siloxane groups of APTES hydrolyze to form Si-OH, which condenses with the metal hydroxyl groups on the coating surface to form Si-OM covalent bonds with a bond energy of 400-450 kJ / mol; POSS nanocages (1-2 nm in diameter) are embedded in the epoxy cross-linked network, acting as physical cross-linking points to increase the resin modulus by 30-40%, while inhibiting crack propagation through steric hindrance effects.

[0043] Preferably, the parameters of the plasma activation treatment in step (a) are: Gas ratio Ar / O2 = 3.8:1-4.2:1, total flow rate 80-100 sccm; RF power 450-550W, treatment time 14-16min, vacuum degree 8×10 -3 ~1.2×10 -2 Pa.

[0044] The volume ratio of Ar to O2 is strictly limited to the range of 3.8:1 to 4.2:1. This ratio is optimized and determined by in-situ monitoring of plasma emission spectroscopy. Ar, as an inert gas, dominates the physical etching process. Its high-energy ions (energy 50 to 100 eV) bombard the surface of the carbon fiber and strip off the surface amorphous carbon impurities; O2, as a reactive gas, participates in chemical modification, and its dissociation generates oxygen free radicals (O*, O2 + ) reacts with the carbon surface to produce carboxyl (-COOH) and hydroxyl (-OH). When the Ar ratio is lower than 3.8:1, excessive O2 causes excessive oxidation of the carbon fiber body (oxidation depth> 50nm), destroying the mechanical properties of the fiber; when the ratio is higher than 4.2:1, insufficient O2 concentration leads to a functional group density (<5 / nm 2 ) cannot meet the chemical bonding requirements of subsequent metal layers.

[0045] The gas ratio of 3.8:1 to 4.2:1 makes the Ar ion etching rate (≈0.8nm / s) and the O2 oxidation rate (≈0.2nm / s) reach a dynamic balance, achieving the best etching depth (≈20nm) and functional group density (8 to 12 / nm). 2 ) is optimized synchronously. At this ratio, Ar + The sputtering effect preferentially removes the weakly bound amorphous carbon on the fiber surface, while O* radicals selectively oxidize the edges of graphite crystallites to form sp 3 The active area is mainly composed of hybrid carbon.

[0046] The total gas flow rate is controlled between 80 and 100 sccm, a range determined through plasma impedance matching experiments. Below 80 sccm, the gas density within the vacuum chamber is insufficient, leading to unstable plasma glow discharge (electron temperature fluctuations >30%). Above 100 sccm, gas turbulence intensifies, weakening the directional ion bombardment effect and significantly reducing surface treatment uniformity (roughness deviation >15%).

[0047] The flow rate of 80-100 sccm stabilizes the gas pressure in the vacuum chamber at 0.5-1.0 Pa. At this time, the thickness of the plasma sheath (≈2 mm) matches the diameter of the carbon fiber (≈5 μm), ensuring that the ion beam is perpendicular to the incident surface. At the same time, the gas residence time (≈0.5 s) at this flow rate matches the RF cycle (13.56 MHz corresponds to a cycle of ≈74 ns), allowing the O2 molecules to fully dissociate in the discharge gap and the free radical concentration to reach a peak value (≈10 15 / cm 3 ).

[0048] The RF power is limited to 450~550W, which is based on the correlation analysis between the plasma electron density (n_e) and the ion energy distribution function (IEDF). When the power is lower than 450W, the electron density is insufficient (n_e<10 10 / cm 3 ), the ion bombardment energy is lower than the carbon fiber surface binding energy threshold (≈3eV), and the etching efficiency decreases; when the power is higher than 550W, high-energy ions (>120eV) penetrate the carbon fiber surface (>50nm) and destroy its graphite microcrystalline structure.

[0049] The radio frequency power of 450-550W keeps the plasma electron density stable at (1.2-1.8)×10 10 / cm 3The ion energy distribution exhibits a bimodal characteristic: a low-energy peak (≈30eV) dominates physical etching, while a high-energy peak (≈80eV) triggers O2 dissociation. Under this energy distribution, nanoscale grooves (depth ≈20nm, width ≈50nm) and a symbiotic structure of functional groups form on the carbon fiber surface. The long axis of the grooves forms an angle of 45° to 60° with the fiber axis, providing a directional growth template for subsequent magnetron sputtering thin films.

[0050] The treatment time was set to 14-16 min, and the vacuum degree was maintained at 8×10 -3 ~1.2×10 -2 When the time is less than 14 min, the surface etching and oxidation reaction have not reached a steady state, and the functional groups are unevenly distributed (CV value>20%). When the time exceeds 16 min, the etching depth enters the carbon fiber reinforcement layer (>30 nm), resulting in a decrease in tensile strength. The vacuum degree is lower than 8×10 -3 Pa, the partial pressure of residual gas (H2O, N2) increases, contaminating the active surface; above 1.2×10 -2 Pa, the mean free path of gas molecules (<5 cm) is shortened and the collision scattering of ion beams is intensified.

[0051] The treatment time of 14 to 16 minutes makes the surface roughness (Ra = 0.8 to 1.2 μm) and the functional group density reach saturation values simultaneously. -3 ~1.2×10 -2 The vacuum degree of Pa ensures a uniform electric field in the plasma sheath (deviation <5%) while suppressing abnormal discharge caused by secondary electron emission. Under these conditions, a hierarchical rough structure consisting of 200-500nm island-like protrusions and nano-grooves is formed on the carbon fiber surface, increasing the specific surface area by about 8 times, providing a dual bonding mechanism of mechanical interlocking and chemical bonding for the metal layer.

[0052] Preferably, the power supply parameters for the three-pulse electroplating in step (d) are: Forward pulse width 0.95~1.05ms, reverse pulse width 0.18~0.22ms, intermittent time 0.48~0.52ms; The current density is dynamically adjusted according to the formula J = 4.0-0.05C_{Al2O3}, where C_{Al2O3} is the concentration of Al2O3 in the plating solution.

[0053] The forward pulse width is limited to 0.95-1.05ms. During this period, nickel ions are reduced and deposited on the cathode surface and nanoparticles are adsorbed. When the pulse width is less than 0.95ms, the electric field is not applied for a long time, and the nanoparticles are stripped by the reverse pulse before they have completed electrophoretic migration to the cathode surface. When the pulse width exceeds 1.05ms, a high current density (>4A / dm 2 ) leads to stress accumulation (>300MPa) in the coating, inducing microcracks.

[0054] The pulse width of 0.95 to 1.05 ms matches the electrophoretic migration time of nano-α-Al2O3 (45 to 55 nm) in the plating solution (≈0.9 ms), ensuring that the particles are fully adsorbed to the cathode surface before the end of the positive pulse; At this pulse width, the critical size of the nickel nucleus (≈5nm) matches the growth rate (≈0.6nm / ms), forming fine equiaxed crystals (grain size <30nm), providing a uniform mosaic matrix for the nanoparticles.

[0055] The reverse pulse width is set to 0.18-0.22ms. During this time, the anodic polarization current briefly dissolves the weakly bonded nickel matrix and nanoparticles on the cathode surface. When the reverse pulse duration is less than 0.18ms, the dissolution only removes free ions adsorbed on the surface and fails to remove particles with weak bonding strength. When the reverse pulse duration exceeds 0.22ms, excessive dissolution can damage the strongly bonded nickel-particle interface.

[0056] Reverse current density (≈8A / dm 2 ) causes the nickel matrix to undergo anodic dissolution (rate ≈1.2μm / min), but the nano-α-Al2O3 (insulating) and silicon carbide whiskers (semiconducting) are retained due to their electrochemical inertness, achieving precise removal of weakly bound particles; The instantaneous high temperature of the reverse pulse (locally >80°C) promotes solid-state diffusion between the nickel matrix and the particle interface, forming a Ni-O-Al transition layer with a thickness of about 2 to 5 nm, thereby improving the interface bonding strength.

[0057] The pause time is strictly controlled between 0.48 and 0.52 ms. During this period, the current returns to zero and the ion concentration gradient in the plating solution is partially restored through diffusion. When the pause time is less than 0.48 ms, the nickel ion concentration in the boundary layer of the plating solution is not fully replenished, resulting in fluctuations in the subsequent pulse deposition rate (deviation >15%). When it exceeds 0.52 ms, the nanoparticles are separated from the cathode surface due to Brownian motion, reducing the effective adsorption rate.

[0058] During the intermittent period, the diffusion flux between the main area of the plating solution (concentration C0) and the nickel-poor area on the cathode surface (concentration C1≈0.2C0) reaches equilibrium, and the nickel ion concentration recovers to above 0.8C0, ensuring a stable deposition rate in the next pulse cycle. During the intermission period, the cathode rotation (28-32 rpm) is synchronously turned on to use centrifugal force (≈5g) to throw the agglomerated particles away from the boundary layer to maintain the particle dispersion stability in the plating solution (particle size deviation <10%).

[0059] The current density is adjusted in real time using the formula J = 4.0-0.05C_{Al2O3}, where C_{Al2O3} is the instantaneous concentration of nano-α-Al2O3 in the plating solution (in g / L). This formula dynamically adjusts the J value by monitoring the Al2O3 concentration through online laser-induced breakdown spectroscopy (LIBS).

[0060] When the Al2O3 concentration increases (from 80g / L to 82g / L), the current density decreases by 0.1A / dm 2 , inhibiting the "shielding effect" caused by excessive particles (i.e. particles covering the cathode surface hindering nickel deposition); The negative correlation between J and C_{Al2O3} maintains a constant ratio (≈4:1) between the nickel substrate deposition rate (V_Ni≈0.8μm / min) and the particle embedding rate (V_p≈0.2μm / min), ensuring that the volume fraction of the nanophase in the coating remains stable at 18-22%.

[0061] Preferably, a closed-loop control system is provided in step (d), comprising: Laser-induced breakdown spectroscopy (LIBS) monitors the Ni and Al2O3 concentrations in the plating solution in real time, triggering refilling when the deviation exceeds ±1.5wt%; a white light interferometer measures the coating thickness online with an accuracy of ±0.4 to 0.6μm.

[0062] Laser induced breakdown spectroscopy (LIBS) was used to investigate the Ni 2+ The Al2O3 concentration is monitored online, and the automatic refill system is triggered when the concentration deviation exceeds ±1.5wt%. The LIBS detection frequency is set to once every 30 seconds, the laser pulse energy is 80-100mJ, and the spectrum acquisition delay is 1.5-2.5μs.

[0063] A high-energy laser (wavelength 1064 nm) penetrates a micro-region (diameter ≈ 100 μm) of the plating solution to generate plasma. The concentration is inverted by the intensity ratio of the Ni characteristic spectral lines (352.4 nm, 361.9 nm) and the Al characteristic spectral lines (394.4 nm, 396.1 nm), with a signal-to-noise ratio (SNR) > 50 dB. The ±1.5wt% deviation threshold is set based on a steady-state mass transfer model for the plating bath. When the Al2O3 concentration exceeds the limit, the feeding system precisely adds nanoparticle slurry according to the formula Δm = 0.12V·ΔC (V is the volume of the plating bath), ensuring that the particle volume fraction fluctuates within ±0.8%. The LIBS optical path is coaxially arranged with the plating bath circulation pipeline, utilizing a stable liquid film in a laminar flow state (Reynolds number <2000) to eliminate bubble scattering interference, achieving a detection accuracy of ±0.3wt%.

[0064] The coating thickness was measured non-contact using a vertical scanning white light interferometer (VSI mode) with a scanning frequency of 10 Hz and an axial resolution of 0.4 to 0.6 μm. The interference fringes were analyzed using a fast Fourier transform (FFT) and the absolute thickness was calculated based on the coating's refractive index (n = 1.35 to 1.45).

[0065] The white light source (wavelength range 480-680nm) is split by a Michelson interferometer, and the precise positioning of the multi-wavelength interference envelope peak (error <λ / 20) is used to break through the 2π ambiguity limit of traditional monochromatic light interferometry. Phase noise caused by coating surface fluctuations (Ra < 0.1 μm) is eliminated by Zernike polynomial fitting to ensure thickness measurement accuracy on a moving cathode (linear speed ≈ 5 mm / s); The thickness data was synchronized with the LIBS concentration signal (time difference < 0.1s), and the plating speed (μm / min) and Ni 2+ A linear relationship model of concentration for predictive feeding control.

[0066] Preferably, during vacuum curing in step (e), the resin film thickness is controlled to be 0.48 to 0.52 mm, and the atomization pressure is 0.28 to 0.32 MPa.

[0067] The film thickness is strictly limited to the range of 0.48 to 0.52 mm, determined through experiments comparing the surface roughness of the nanocomposite coating (Ra = 0.8 to 1.2 μm) with the resin viscosity (2000 to 2500 mPa·s). When the film thickness is less than 0.48 mm, the resin cannot completely fill the micropores and whisker gaps (pore diameter 1 to 3 μm) on the coating surface, resulting in a weakened mechanical interlocking effect at the interface. When the film thickness exceeds 0.52 mm, the curing shrinkage stress (≈15 MPa) within the resin layer exceeds the coating bonding strength threshold, causing the risk of edge warping.

[0068] The wet film thickness of 0.48-0.52mm is affected by the synergistic effect of gravity and capillary force during the vacuum curing stage (the first stage). The resin penetrates into the pores of the coating to a depth of 20-50μm, forming a "barb-like" anchoring structure. The film thickness matches the POSS nanoparticle content (4.8-5.2 wt%) in the resin, and the rigid cage structure of POSS acts as a stress dispersion point, reducing residual stress after curing by 40-50%. The film thickness control ensures that the concentration gradient of the APTES silane coupling agent at the interface is optimal, and the condensation reaction coverage of the silanol (Si-OH) and the metal hydroxyl (M-OH) of the coating is ≥85%.

[0069] The atomization pressure is set between 0.28 and 0.32 MPa, a range determined by calculating the relationship between droplet size distribution and surface energy. When the pressure is below 0.28 MPa, the resin droplet diameter (D50 > 50 μm) is too large, and after impacting the coating surface, it does not spread sufficiently, forming localized accumulation (thickness deviation > 15%). When the pressure is above 0.32 MPa, the droplets are too fine (D50 < 10 μm), causing the solvent to evaporate too quickly, prematurely gelling the resin, and hindering interfacial penetration.

[0070] Under a pressure of 0.28-0.32 MPa, the atomized droplet D50 = 20-30 μm, its Weber number (We≈5-10) matches the wettability of the coating surface (contact angle ≈60°), and a uniform liquid film is formed after impact (thickness fluctuation <±5%); ②Solvent evaporation control: droplet flight time (≈0.2s) and solvent (acetone) evaporation rate (≈0.1g / m 2 s) matching, ensuring that the solid content of the resin is 65-70% when it reaches the surface of the coating, retaining sufficient fluidity to achieve pore filling; ③Nanodispersion enhancement: atomization shear force (≈10 4 s -1 ) to make the POSS particles (particle size 1-2nm) uniformly dispersed in the resin (agglomerate size <50nm), avoiding stress concentration caused by excessive local concentration.

[0071] In summary, the present invention includes at least one of the following beneficial technical effects: 1. This invention achieves cross-scale synergistic reinforcement of the metal matrix, ceramic layer, and resin layer through a gradient composite interface structure design. The dual buffering mechanism of the porous ceramic layer and nanocomposite coating effectively disperses interfacial stress concentration. The formation of a nano-crosslinked network in the hybrid resin further enhances the synergistic effect of chemical bonding and physical anchoring, enabling the material to maintain excellent interfacial stability even in a thermomechanically coupled environment.

[0072] 2. This invention innovatively constructs a "soft-hard-soft" cross-scale transition structure based on a gradient matching strategy for thermal expansion coefficient and elastic modulus. This design achieves a step-by-step release of thermal stress through the synergistic energy dissipation mechanism of amorphous phase plastic deformation and nanocrystalline phase slip, resolving the problem of interfacial cracking caused by material thermal expansion differences during high-temperature service.

[0073] 3. This invention utilizes a dynamic electroplating process and multi-physics field coupling control technology to overcome the technical challenges of nanoparticle agglomeration and random whisker orientation. Through the synergistic effect of the electric and fluid fields, the migration paths and adsorption sites of the nanophase are precisely controlled, resulting in a composite reinforcement structure of diffusion strengthening and crack bridging within the coating, significantly enhancing the material's erosion and fatigue resistance.

[0074] 4. This invention deeply integrates plasma activation, micro-arc oxidation, and closed-loop control technologies to form a controllable preparation method that integrates composition, structure, and performance. Through dynamic feedback of process parameters and coordinated optimization of multiple steps, this system ensures precise formation and stable performance of cross-scale structures, providing a reliable solution for the large-scale production of high-performance composite materials under complex operating conditions. BRIEF DESCRIPTION OF THE DRAWINGS

[0075] Figure 1 Schematic diagram of the method of the present invention. DETAILED DESCRIPTION

[0076] The following is combined with Figure 1 , the present invention is described in further detail.

[0077] Example 1 Step (a) plasma activation treatment: Gas ratio: Ar / O2=4:1 (volume ratio); Total flow rate: 90 sccm; RF power: 500W; Processing time: 15 minutes; Vacuum degree: 1.0×10 -2 Pa.

[0078] Step (b) magnetron sputtering Ni-Ti transition layer: Target composition: Ni 50 Ti 50 (atomic ratio); Sputtering power: 300W; Substrate temperature: 200°C; Deposition rate: 0.5 nm / s; Final thickness: 20 μm.

[0079] Step (c) micro-arc oxidation treatment: Electrolyte composition: Na2SiO310g / L+KOH5g / L; Pulse voltage: 450V; Frequency: 500Hz; Processing time: 10 minutes; Ceramic layer thickness: 10 μm; Porosity: 35%.

[0080] Step (d) three-pulse electroplating nanocomposite coating: Plating solution composition: nickel sulfamate 380g / L + α-Al2O3 80g / L + SiC whisker 15g / L; forward pulse width: 1.0ms; Reverse pulse width: 0.2ms; Interval time: 0.5ms; Current density: According to the formula J = 4.0-0.05×80 = 3.6A / dm 2 Dynamic adjustment.

[0081] Step (e) Preparation of epoxy-silane hybrid protective layer: Resin ratio: E51 epoxy resin 95wt% + APTES 8wt% + POSS 5wt%; Wet film thickness: 0.50mm; Atomization pressure: 0.30MPa; Curing procedure: 80℃ / 2h→120℃ / 1h→150℃ / 0.5h.

[0082] Example 2 Step (a) plasma activation treatment: Gas ratio: Ar / O2=3.8:1; Total flow rate: 80 sccm; RF power: 450W; Processing time: 14 minutes; Vacuum degree: 8×10 -3 Pa.

[0083] Step (b) magnetron sputtering Ni-Ti transition layer: Target composition: Ni 50 Ti 50 ; Sputtering power: 280W; Substrate temperature: 180°C; Deposition rate: 0.45 nm / s; Final thickness: 18 μm.

[0084] Step (c) micro-arc oxidation treatment: Electrolyte composition: Na2SiO39.5g / L+KOH4.8g / L; Pulse voltage: 440V; Frequency: 480Hz; Processing time: 9 minutes; Ceramic layer thickness: 8μm; Porosity: 30%.

[0085] Step (d) three-pulse electroplating nanocomposite coating: Plating solution composition: nickel sulfamate 375g / L + α-Al2O3 78g / L + SiC whisker 14g / L; forward pulse width: 0.95ms; Reverse pulse width: 0.18ms; Interval time: 0.48ms; Current density: According to the formula J = 4.0-0.05×78 = 4.0-3.9 = 0.1A / dm 2 .

[0086] Step (e) Preparation of epoxy-silane hybrid protective layer: Resin ratio: E51 epoxy resin 94wt% + APTES 7.5wt% + POSS 4.8wt%; wet film thickness: 0.48mm; Atomization pressure: 0.28MPa; Curing procedure: 78℃ / 2.1h→118℃ / 1.05h→148℃ / 0.52h.

[0087] Example 3 Step (a) plasma activation treatment: Gas ratio: Ar / O2=4.2:1; Total flow rate: 100 sccm; RF power: 550W; Processing time: 16 minutes; Vacuum degree: 1.2×10 -2 Pa.

[0088] Step (b) magnetron sputtering Ni-Ti transition layer: Target composition: Ni 50 Ti 50 ; Sputtering power: 320W; Substrate temperature: 220℃; Deposition rate: 0.55 nm / s; Final thickness: 22 μm.

[0089] Step (c) micro-arc oxidation treatment: Electrolyte composition: Na2SiO310.5g / L+KOH5.2g / L; Pulse voltage: 460V; Frequency: 520Hz; Processing time: 11 minutes; Ceramic layer thickness: 12μm; Porosity: 40%.

[0090] Step (d) three-pulse electroplating nanocomposite coating: Plating solution composition: nickel sulfamate 385g / L + α-Al2O3 82g / L + SiC whisker 16g / L; forward pulse width: 1.05ms; Reverse pulse width: 0.22ms; Interval time: 0.52ms; Current density: According to the formula J = 4.0-0.05×82 = 4.0-4.1 = -0.1A / dm 2 .

[0091] Step (e) Preparation of epoxy-silane hybrid protective layer: Resin ratio: E51 epoxy resin 96wt% + APTES 8.5wt% + POSS 5.2wt%; wet film thickness: 0.52mm; Atomization pressure: 0.32MPa.

[0092] Curing procedure: 82℃ / 1.9h→122℃ / 0.95h→152℃ / 0.48h.

[0093] Comparative Example 1 Compared with Example 1, the difference is: In step (a) of the plasma activation treatment, the gas ratio was changed to Ar / O2=5:1, and the treatment time was shortened to 8 minutes. The remaining steps and parameters remained the same.

[0094] Comparative Example 2 Compared with Example 1, the difference is: Step (b) magnetron sputtering of the Ni-Ti transition layer was omitted, and step (c) micro-arc oxidation treatment was directly performed on the surface of the activated carbon fiber. The remaining steps and parameters were the same.

[0095] Comparative Example 3 Compared with Example 1, the difference is: In step (c) micro-arc oxidation treatment, the pulse voltage was increased to 500 V, and the electrolyte was changed to a single component Na2SiO3 (15 g / L). The remaining steps and parameters were the same.

[0096] Comparative Example 4 Compared with Example 1, the difference is: Step (d) The three-pulse electroplating was changed to a single-pulse electroplating (forward pulse width 1.0 ms, no reverse pulse and rest time), and the current density was fixed at 3.6 A / dm 2 , the rest of the steps and parameters are the same.

[0097] Comparative Example 5 Compared with Example 1, the difference is: Cancel the closed loop control system, Ni plating solution 2+ The Al2O3 concentration was manually sampled every 2 hours, and the current density was fixed at 3.6A / dm 2 , the rest of the steps and parameters are the same.

[0098] Comparative Example 6 Compared with Example 1, the difference is: In step (e), the POSS component is omitted from the epoxy-silane hybrid resin, and only E51 epoxy resin and APTES (the ratio is adjusted to E51 95 wt % + APTES 10 wt %) are used, and the remaining steps and parameters are all the same.

[0099] Comparative Example 7 Compared with Example 1, the difference is: In step (e), the curing procedure was changed to single-stage curing (150° C. / 3.5 h), and the vacuum environment and staged temperature increase were eliminated. The remaining steps and parameters remained the same.

[0100] Test Example 1: Interface Bonding Strength and Thermal Stress Stability Experimental Description Experimental procedures Sample preparation: The coating samples of Examples 1-3 and Comparative Examples 1, 2, and 7 were cut into a size of 25 mm×10 mm×3 mm (length×width×thickness).

[0101] The sample surface was ultrasonically cleaned with acetone (20 min), and after drying, a 0.1 mm thick gold film (for SEM conductive layer) was sprayed on it.

[0102] Lap shear strength test: Equipment: Universal material testing machine (Instron 5967), high temperature fixture (RT ~ 300℃) condition: The overlap area is 5mm×5mm, and the epoxy structural adhesive (3MDP460) is used for bonding; The heating rate was 5°C / min, and the samples were loaded until fracture at 25°C, 150°C, and 300°C, respectively, with a loading rate of 1 mm / min.

[0103] Data recording: For each group of 3 parallel samples, take the maximum breaking load.

[0104] Liquid nitrogen-high temperature thermal shock experiment: Equipment: Liquid nitrogen tank (-196°C), muffle furnace (300°C), cycle counter process: The sample was immersed in liquid nitrogen for 5 minutes, then transferred to a muffle furnace at 300°C and kept for 5 minutes, which was defined as one cycle. After repeating 100 times, the percentage of coating peeling area was calculated using ImageJ software.

[0105] Interface micromorphology analysis Equipment: Field emission scanning electron microscope (FEINovaNanoSEM450) Sample preparation: Cutting along the interface direction, argon ion polishing (GatanPECSII); The crack propagation path and pore distribution were observed in low vacuum mode (10Pa).

[0106] The experimental data are shown in Table 1 below: Table 1 - Comparison of interface bonding strength and peeling area after thermal shock The precise control of the Ar / O2 gas ratio and treatment time during plasma activation directly determines the morphology of the nanogrooves and the functional group density on the carbon fiber surface. In Comparative Example 1, due to the unbalanced gas ratio (Ar / O2 = 5:1) and insufficient treatment time, the surface oxygen content was only increased to 7.2 at%, which is significantly lower than the 12.3 at% in Example 1. + The bombardment ratio leads to excessive surface etching, insufficient exposure of the graphite crystallite edges, and reduced nucleation site density in the subsequent Ni-Ti transition layer. Microscopic morphology reveals continuous delamination (>200μm) at the interface of Comparative Example 1, confirming the degradation of the mechanical anchoring effect caused by plasma parameter mismatch, which is consistent with the "dynamic balance between physical etching and chemical modification" logic in the previous mechanism.

[0107] The amorphous-nanocrystalline composite structure of the Ni-Ti transition layer plays a key role in thermal stress buffering. In Comparative Example 2, the transition layer is omitted, and the carbon fiber (CTE≈0.5×10 -6 / ℃) and nanocomposite coating (CTE≈14×10 -6 / ℃) in direct contact with each other, the interface shear strength dropped sharply to 6.4MPa at 300℃, which is only 15% of that in Example 1. After the thermal shock test, the coating peeled off as much as 89.5%, indicating that the drastic CTE difference caused the interface shear stress concentration. This is consistent with the previous "thermal expansion coefficient gradient design" mechanism: the CTE of the Ni-Ti layer (11×10 -6 / ℃) as an intermediate transition phase, the stress is dissipated synergistically through the plastic deformation of the amorphous matrix and the dislocation slip of the nanocrystalline phase, thus avoiding the brittle peeling of the ceramic layer.

[0108] Phased vacuum curing process optimizes the rheological behavior and chemical bonding process of resin by temperature-pressure sequential design.Comparative example 7 adopts single-stage high-temperature curing, and resin rapidly gels at 150 ℃, and APTES silane coupling agent does not fully penetrate into ceramic layer pore interior (penetration depth is only 8 μm, and embodiment 1 is 30 μm), causes Si-O-Al bonding coverage less than 45%.Meanwhile, POSS nanoparticles are localized and agglomerated (size>200nm) too fast in the initial stage of solidification because of solvent volatilization, cause resin layer internal stress concentration, and interface microcrack density reaches 10 / mm.This result verifies the synergistic necessity of " capillary force driven penetration " and " POSS stress dispersion " in mechanism--80 ℃ of pre-curing stages maintain resin fluidity to complete pore filling, and 150 ℃ of final curing stages then release shrinkage stress by slowly cross-linking, realize interface chemical-mechanical joint strengthening.

[0109] Test Example 2: Extreme Environment Service Performance Experiment Description Experimental procedures Sample preparation: Take the samples of Examples 1-3 and Comparative Examples 3 and 6, with dimensions of 50 mm × 50 mm × 3 mm (length × width × thickness); The surface was lightly polished with 600-grit sandpaper (to remove the residue of the sprayed gold film), wiped with ethanol, and dried for 24 hours.

[0110] Sand and dust erosion test: Equipment: Gas-solid two-phase flow erosion test machine (modified according to ASTM G76 standard) condition: Erosion medium: quartz sand (particle size 150±20μm, Mohs hardness 7); Jet speed: 80±5m / s; Erosion angle: 90° (vertical impact); Test time: 30 minutes (accumulated sand volume is about 2.5 kg).

[0111] Data recording: An electronic balance (accuracy 0.1 mg) was used to measure the mass difference before and after erosion and calculate the mass loss rate.

[0112] Vacuum ultraviolet radiation aging: Equipment: Space environment simulation box (vacuum degree 10 -5 Pa, xenon lamp ultraviolet band 200-400nm) process: The sample was fixed on a rotating stand (5 rpm) with a UV intensity of 5.0 ± 0.3 kW / m 2 ; After continuous irradiation for 500 h, the surface powdering degree was observed using a colorimeter (ΔE value) and SEM.

[0113] Salt spray corrosion test: Equipment: Circulating salt spray chamber (Q-FOGCCT1100) condition: Solution: 5wt% NaCl, pH=6.5-7.2; Temperature: 35±1℃; Spray / dry cycle: spray 2 h → dry 4 h → repeat 30 times (720 h); Analysis: Laser confocal microscopy was used to measure the maximum depth of the corrosion pits and the density (number / mm 2 ).

[0114] The experimental data are shown in Table 2 below: Table 2-Comparison of service performance in extreme environments The precise control of pulse voltage and electrolyte composition during micro-arc oxidation determines the bimodal distribution characteristics of the ceramic layer pores. In Comparative Example 3, the voltage was increased to 500V and the KOH component was eliminated, resulting in a single discharge pulse energy that was too high (>3000J / cm 2 ), the diameter of the molten pool expanded to more than 15μm, and the main pore size exceeded the limit to 5-8μm. The coarse pores weakened the composite structural effect of "large pore diversion + small pore anchoring". When the quartz sand impacted, the stress concentrated on the weak area of the pore wall, causing collapse damage (diameter>50μm). This is consistent with the control logic of "high-frequency pulse limiting the molten pool size" in the early mechanism - the amorphous matrix (70%) formed by the rapid solidification of the molten pool at a frequency of 500Hz can constrain the pore expansion, and SiO3 2- The decomposed nano-SiO2 (≈50nm) fills the sub-pore gaps and jointly maintains the mechanical stability of the pore structure.

[0115] The nano-enhancement effect of POSS in epoxy-silane hybrid resin significantly affects the environmental aging resistance. Comparative Example 6 omits the POSS component. During the resin curing process, the APTES silane coupling agent is only combined with the ceramic layer surface through the Si-O-Al bond, and fails to form a covalent cross-linked network of POSS and epoxy resin. Under ultraviolet irradiation, the molecular segment of pure epoxy resin breaks due to photooxidation reaction, and the surface powdered area (coverage ≈ 30%) is accompanied by microcrack initiation (density 15 / mm), and the ΔE value increases to more than 12.3. This confirms the "steric hindrance effect of POSS nanoparticles" in the mechanism - its cage structure suppresses crack propagation through steric hindrance, and the Si-O-Si network promotes the ultraviolet shielding efficiency of the resin layer, so that the ΔE value of Example 1 is stabilized at 3.2-3.5.

[0116] The orientation distribution of SiC whiskers and the uniform dispersion of α-Al2O3 particles in the nanocomposite coating are the key factors in resisting sand dust erosion and salt spray corrosion. Example 3 Dynamic current regulation (J = 4.0-0.05×82 = 3.6A / dm 2 ), so that the SiC whiskers are oriented along the direction of the electric field (orientation angle deviation <8°), and the exposed height of the whiskers is controlled within 50nm, forming dense mechanical interlocking points. In Example 2, due to the lower limit of the Al2O3 concentration in the plating solution (78g / L), the spatial competitive adsorption between the whiskers and the particles causes the volume fraction of the nanophase in the local area to drop to 17%, and the depth of the salt spray corrosion pit increases to 12.8μm. This result verifies the synergistic mechanism of "dielectrophoresis effect drives whisker orientation" and "dynamic current compensates for concentration fluctuations" - the forward pulse induces whisker orientation, and the reverse pulse dissolves the weakly bonded nickel matrix, ensuring that the coating has both high hardness (HV0.1=520) and corrosion resistance (corrosion rate <0.01mm / year).

[0117] Test Example 3: Experimental description of coating structure and composition uniformity Experimental procedures Sample preparation: The coating samples of Examples 1-3 and Comparative Examples 4 and 5 were cut into 10 mm × 10 mm slices; The sample was thinned to 80-100 nm (for TEM observation) using an ion thinning instrument (Gatan 691); The unthinned specimens were mounted with epoxy resin and mechanically polished to Ra < 0.05 μm (for EDS analysis).

[0118] Transmission electron microscopy (TEM) nanostructure analysis: Equipment: Field emission transmission electron microscope (JEOL JEM-ARM200F) condition: Accelerating voltage 200 kV, STEM-HAADF mode; Randomly select 5 fields of view (5 μm each field of view) 2 ), and statistics were made on the spacing between α-Al2O3 particles and the orientation angle of SiC whiskers (the angle with the coating surface).

[0119] Composition line scanning and surface distribution analysis: Equipment: Energy dispersive spectrometer (EDS, Oxford X-Max N150) process: Line scanning was performed along the coating thickness direction (from surface to substrate) with a step size of 1 μm and a total length of 20 μm; Surface scanning analysis (100 μm × 100 μm) was used to obtain the uniformity of Ni and Al element distribution (calculate the coefficient of variation CV value).

[0120] Plating rate stability monitoring: Equipment: Electroplating process data acquisition system (National Instruments cDAQ-9188) method: Real-time recording of electroplating current density fluctuations (sampling rate 10 Hz) and calculation of standard deviation (σ); The coating thickness was measured every 10 min (white light interferometer, Zygo New View 9000), and the slope of the thickness-time curve was fitted.

[0121] The experimental data are shown in Table 3 below: Table 3 - Comparison of coating structure and composition uniformity The strict timing matching of the forward / reverse pulse width and the intermittent time of three-pulse electroplating is the decisive factor for the uniform dispersion of nanoparticles and the oriented arrangement of whiskers. After the single-pulse mode was adopted in Comparative Example 4, the continuous loading of the forward current caused the α-Al2O3 particles to be excessively adsorbed on the cathode surface, forming local agglomerations (spacing>150nm), and the absence of the reverse pulse prevented the weakly bound SiC whiskers from being peeled off, and the orientation angle deviation expanded to 34.6°. This is directly related to the "forward adsorption-reverse screening" regulation mechanism in the early mechanism - the forward pulse width (1.0ms) provides sufficient electric field to drive the dielectrophoretic force, prompting the nanoparticles to migrate to the cathode; the reverse pulse width (0.2ms) dissolves the loosely adsorbed particles through anodic polarization, retaining only the strongly bound whiskers (binding energy>1.5eV), thereby achieving spatially ordered distribution of the nanophase.

[0122] The real-time feedback capability of the dynamic closed-loop control system is crucial to the stability of the plating solution composition. In Example 5, due to the cancellation of LIBS online monitoring and thickness feedback, the artificial feeding delay caused the instantaneous fluctuation of Al2O3 concentration to reach ±4.5g / L, which caused the current density regulation to lag (response time>30s). An obvious Ni concentration gradient (1.32at% / μm) appeared in the thickness direction of the coating, and the CV value of the Al element surface distribution rose to 12.5%, which was much higher than 4.7% in Example 1. This result verifies the effectiveness of the "concentration-current linkage model" in the mechanism - the formula J=4.0-0.05C controls the fluctuation of the Al2O3 volume fraction within ±0.8% by compensating for the particle shielding effect in real time, ensuring the linear correlation between the uniformity of the nanophase distribution and the mechanical properties of the coating.

[0123] The synergistic effect of high-frequency pulsed electric field and ultrasonic flow field optimizes the orientation consistency of whiskers. Example 3 breaks the concentration polarization of the plating solution boundary layer by coupling of cathode rotation (32rpm) and ultrasonic wave (40kHz), so that SiC whiskers are directional deflected along the electric field lines under the action of dielectrophoretic force (orientation angle <6.3 °). And comparative example 4 lacks dynamic flow field disturbance due to single-pulse electroplating, and whiskers are randomly distributed under the eddy current effect (orientation angle>30 °), which weakens its crack bridging function. This phenomenon is consistent with the logic of "flow field-electric field coupling regulation" described in the mechanism - ultrasonic cavitation effect produces micron-level turbulence (Reynolds number>5000), forcing whiskers to continuously adjust their spatial posture during migration, and finally achieves high-precision alignment with the principal stress direction (deviation angle<5 °), significantly improving the coating's ability to resist crack growth.

[0124] While embodiments of the present invention have been shown and described, it will be appreciated by those skilled in the art that various changes, modifications, substitutions, and variations may be made to these embodiments without departing from the principles and spirit of the invention, and that the scope of the invention is defined by the appended claims and their equivalents.

Claims

1. A modified carbon fiber material for a drone, characterized in that: The following structural layers are included: carbon fiber matrix; The Ni-Ti transition layer deposited on the substrate surface has a thickness of 18 to 22 μm and a mass ratio of Ni to Ti of 49.5:50.5 to 50.5:49.5; The nano-composite coating covering the transition layer comprises a nickel sulfamate matrix and nano-α-Al2O3 and silicon carbide whiskers dispersed therein, wherein the Al2O3 content is 78-82 g / L and the silicon carbide whisker content is 14-16 g / L; The epoxy-silane hybrid protective layer coated on the surface of the composite coating comprises 94-96 wt% of epoxy resin E51, 7.5-8.5 wt% of APTES and 4.8-5.2 wt% of POSS.

2. The modified carbon fiber material for a drone according to claim 1, characterized in that: The Ni-Ti transition layer is prepared by magnetron sputtering, with a sputtering power of 280-320W, a substrate temperature of 180-220°C, and a deposition rate of 0.45-0.55nm / s.

3. The modified carbon fiber material for a drone according to claim 1, characterized in that: In the nano-composite coating, the particle size of nano-α-Al 3 O 3 is 45-55 nm, the diameter of the silicon carbide whisker is 180-220 nm, and the coating thickness is 78-82 μm.

4. The modified carbon fiber material for a drone according to claim 1, characterized in that: The curing process of the epoxy-silane hybrid protective layer includes three stages: Stage 1: Curing at 78-82°C for 1.9-2.1h, vacuum degree -0.094-0.096MPa; The second stage: curing at 118-122°C for 0.95-1.05h, vacuum degree -0.094-0.096MPa; The third stage: curing at 148-152°C for 0.48-0.52h, normal pressure environment.

5. The modified carbon fiber material for a drone according to claim 1, characterized in that: A porous ceramic layer generated by micro-arc oxidation is provided between the nano-composite coating and the Ni-Ti transition layer, with a porosity of 30-40% and a thickness of 8-12 μm.

6. A processing technology based on the modified carbon fiber material according to any one of claims 1 to 5, characterized in that: The following steps are involved: (a) Plasma activation treatment of carbon fiber matrix; (b) Deposition of a Ni-Ti transition layer on the activated surface by magnetron sputtering; (c) performing micro-arc oxidation treatment on the surface of the transition layer to form a porous ceramic layer; (d) Preparation of nanocomposite coating on porous ceramic layer using triple-pulse electroplating process; (e) Spraying epoxy-silane hybrid resin on the surface of the composite coating and vacuum curing.

7. The processing technology of modified carbon fiber material according to claim 6, characterized in that: The parameters of the plasma activation treatment in step (a) are: Gas ratio Ar / O2=3.8:1~4.2:1, total flow rate 80~100sccm; RF power 450-550W, treatment time 14-16min, vacuum degree 8×10 -3 ~1.2×10 -2 Pa.

8. The processing technology of modified carbon fiber material according to claim 6, characterized in that: The power supply parameters for the three-pulse electroplating in step (d) are: Forward pulse width 0.95~1.05ms, reverse pulse width 0.18~0.22ms, intermittent time 0.48~0.52ms; The current density is dynamically adjusted according to the formula J=4.0-0.05C_{Al2O3}, where C_{Al2O3} is the Al2O3 concentration in the plating solution.

9. The processing technology of modified carbon fiber material according to claim 6, characterized in that: In step (d), a closed-loop control system is set up, including: Laser-induced breakdown spectroscopy (LIBS) monitors the Ni and Al2O3 concentrations in the plating solution in real time, triggering refilling when the deviation exceeds ±1.5wt%; The white light interferometer measures the coating thickness online with an accuracy of ±0.4 to 0.6 μm.

10. The processing technology of modified carbon fiber material according to claim 6, characterized in that: During vacuum curing in step (e), the resin film thickness is controlled to be 0.48-0.52 mm, and the atomization pressure is controlled to be 0.28-0.32 MPa.

Images