A cobalt-chromium alloy laser cladding baked porcelain denture base and a preparation method thereof

Abstract

The present application relates to the technical field of dental restoration materials, and discloses a cobalt-chromium alloy laser cladding baked porcelain denture base and a preparation method thereof. The base is made of cobalt-chromium-molybdenum alloy powder, silicon carbide, niobium pentoxide and tantalum powder through selective laser melting and heat treatment, has dispersedly distributed tantalum carbide phase inside, and has a niobium pentoxide layer on the surface. The preparation method comprises the following steps: mixing the powders according to the proportion; performing focused scanning melting and off-focus scanning slagging on the single-layer spread powder in sequence to form layer by layer; performing heat treatment in a low dew point reducing atmosphere to generate a niobium pentoxide transition layer on the surface of the base; performing isothermal vacuum degassing before the porcelain powder is heated to the glass transition temperature, and then heating to complete the fused porcelain baking. According to the present application, the tantalum carbide is generated in situ to avoid the formation of chromium-poor areas in the matrix, and the niobium pentoxide transition layer is combined with the porcelain layer to improve the gold-porcelain bonding strength, the density, the corrosion resistance and the thermal shock fatigue life of the base.

Description

Technical Field

[0001] This invention relates to the field of dental restorative materials technology, specifically to a cobalt-chromium alloy laser-clad porcelain denture substrate and its preparation method. Background Technology

[0002] Cobalt-chromium alloys possess certain mechanical strength and biocompatibility, and are commonly used as the metal substrate material for porcelain dentures. Selective laser melting technology can achieve digital customization of denture substrates, but during processing, the metal powder undergoes rapid melting and solidification, as well as multiple thermal cycles. These specific metallurgical conditions cause the carbon elements inside the cobalt-chromium alloy to easily react with the chromium elements in the matrix solid solution, resulting in the precipitation of a large amount of network-distributed chromium carbide phase at the grain boundaries.

[0003] The precipitation of chromium carbide consumes chromium around the grain boundaries, forming chromium-depleted regions at the microscopic level. The presence of these chromium-depleted regions disrupts the integrity of the passivation film on the alloy surface, reducing the material's resistance to electrochemical corrosion in the oral saliva environment and increasing the likelihood of metal ions precipitating into human tissue. Simultaneously, the continuously distributed brittle chromium carbide phase hinders dislocation slip during metal deformation under stress, inducing grain boundary embrittlement. This results in a reduced elongation at break after forming, making the material prone to brittle fracture under chewing stress. Summary of the Invention

[0004] To address the shortcomings of existing technologies, this invention provides a cobalt-chromium alloy laser cladding porcelain denture substrate and its preparation method, which solves the problems of reduced plasticity and localized corrosion resistance of the substrate alloy in existing technologies.

[0005] To achieve the above objectives, the present invention provides the following technical solution: a cobalt-chromium alloy laser-clad porcelain denture base, comprising the following raw materials in parts by weight: 100 parts of cobalt-chromium-molybdenum alloy powder; 0.05-0.15 parts of silicon carbide; 0.15-0.40 parts of niobium pentoxide; 0.35-0.65 parts of tantalum powder; The substrate contains a tantalum carbide phase that is used to lock free carbon in situ within its internal microstructure, and the surface of the substrate has a niobium pentoxide transition layer that blocks the oxidation of chromium and forms chemical bonds with the ceramic powder.

[0006] Preferably, the cobalt-chromium-molybdenum alloy powder, silicon carbide, niobium pentoxide and tantalum powder are all spherical in shape, and the particle size distribution range of the powder is 15-53 μm.

[0007] Preferably, the substrate has no micron-sized amorphous silica inclusions in its internal metal matrix, and the macroscopic relative density of the substrate is greater than or equal to 99.40%.

[0008] Preferably, the substrate has a room temperature yield strength greater than or equal to 690 MPa and an elongation at break greater than or equal to 11.0%.

[0009] Preferably, the surface of the substrate is further coated with a feldspar dental porcelain layer, wherein the silicon-oxygen tetrahedral glass network at the bottom of the feldspar dental porcelain layer and the niobium pentoxide transition layer on the surface of the substrate have an ion-level interdiffusion diffusion structure.

[0010] A method for preparing a cobalt-chromium alloy laser-clad porcelain denture substrate includes the following steps: S1. Powder preparation: Weigh out cobalt-chromium-molybdenum alloy powder, silicon carbide, niobium pentoxide and tantalum powder according to the weight ratio, and mix them evenly to obtain composite alloy powder; S2. Selective laser melting and forming: The composite alloy powder is introduced into a selective laser melting device, and the first focused scanning melting is performed on the single-layer spread powder. Then, the second defocusing scanning is performed in situ to induce Marangoni convection for rheological slag removal. The metal substrate is stacked layer by layer to obtain the initial product. S3. Controlled atmosphere heat treatment: The metal substrate sample is placed in a reducing mixed gas atmosphere with an extremely low dew point for heat treatment, which triggers thermodynamic selective oxidation and generates a niobium pentoxide transition layer in situ on the substrate surface to obtain a modified metal substrate. S4. Isothermal degassing porcelain: Dental porcelain powder slurry is applied to the surface of the modified metal substrate. During the heating and sintering process, isothermal vacuum degassing is performed before the porcelain powder crosses the glass transition temperature. Then, the temperature is continued to rise to complete the melting and sintering, thus obtaining the final product.

[0011] Preferably, in step S1, silicon carbide, niobium pentoxide, and tantalum powder are pre-treated as follows: the weighed silicon carbide, niobium pentoxide, and tantalum powder are placed together in a planetary ball mill, and anhydrous ethanol is used as the grinding medium. They are pre-mixed and dispersed at a speed of 250-350 rpm for 2-4 hours. After being vacuum dried at 80°C, they are added to a three-dimensional mixer with cobalt-chromium-molybdenum alloy powder for dry mixing.

[0012] Preferably, the specific implementation of step S2 is as follows: The process parameters for the first focusing scan are: laser power 180-220W, scanning speed 800-1000mm / s, and spot diameter 50-80μm; The process parameters for the second defocus scanning are as follows: the laser power is reduced to 80-120W, the scanning speed is reduced to 400-600mm / s, and the spot diameter is expanded to 120-180μm through positive defocus.

[0013] Preferably, in step S3, the reducing mixed atmosphere is composed of argon and hydrogen with a volume fraction of 5-10%, and the dew point of the reducing mixed atmosphere is strictly controlled to be lower than or equal to -40°C; the heating rate of the heat treatment is controlled to be 10-15°C / min, the holding temperature is 900-1000°C, and the holding time is 1-2h.

[0014] Preferably, in step S4, the isothermal vacuum degassing process is specifically implemented as follows: When the ceramic furnace reaches 580℃, stop heating and maintain an isothermal state at 580℃ while simultaneously drawing a vacuum from the furnace cavity until the absolute pressure in the furnace cavity is below 50Pa. The pressure holding and degassing time should be controlled at 10-15 minutes.

[0015] This invention provides a cobalt-chromium alloy laser-clad porcelain denture substrate and its preparation method. It has the following beneficial effects: 1. This invention introduces niobium pentoxide into cobalt-chromium-molybdenum alloy powder and performs heat treatment in an extremely low dew point reducing atmosphere. This method utilizes a thermodynamic selective oxidation mechanism to generate a niobium pentoxide transition layer in situ on the substrate surface. This structure replaces the loose and brittle oxide scale formed by excessive oxidation of chromium in traditional processes. During the subsequent porcelain-forming stage, it can undergo deep chemical dissolution and ion interpenetration with the silica-oxygen network of feldspar porcelain powder, establishing a tight chemical bond at the interface. This effectively conducts and disperses external mechanical shear loads, reducing the risk of porcelain layer spalling and fracture during clinical service.

[0016] 2. This invention improves the electrochemical corrosion resistance of the alloy and limits the precipitation of harmful metal ions. The method simultaneously adds silicon carbide and tantalum powder to the powder raw material. Utilizing the preferential affinity of tantalum for carbon, the free carbon released during the selective laser melting process is locked in situ into a stable tantalum carbide dispersed phase. This microalloying carbon fixation mechanism fundamentally blocks the consumption of chromium in the matrix solid solution by free carbon, avoiding the precipitation of the hard and brittle chromium carbide phase at grain boundaries and the formation of chromium-depleted zones. This not only ensures the density and integrity of the protective passivation layer on the matrix surface and significantly reduces the self-corrosion current density of the material, but also cuts off the intergranular corrosion channels, ensuring the safety of dentures against ion precipitation in the complex saliva environment of the oral cavity.

[0017] 3. In the molding stage, this invention employs a dual-channel scanning process of focusing and defocusing. The expanded light spot excites stable Marangoni convection, causing low-density inclusions such as amorphous silica to float and be removed, ensuring the structural continuity and high density of the metal interior. Furthermore, in the porcelain-baking process, an isothermal vacuum degassing treatment is introduced before the glass transition temperature. This utilizes the unmelted pores of the porcelain powder to fully remove residual gas from the interface, eliminating microscopic closed pores at the metal-ceramic junction. These two processes synergistically eliminate stress concentration sources both internally and at the interface, significantly improving the yield strength and elongation at break of the substrate, and substantially extending the number of thermal fatigue failure cycles the restoration can withstand. Attached Figure Description

[0018] Figure 1 This is a schematic diagram of the method steps of the present invention. Detailed Implementation

[0019] The technical solution of the present invention will now be clearly and completely described with reference to the accompanying drawings. 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.

[0020] To make the objectives, technical solutions, and advantages of this invention clearer, the invention will be further described in detail below with reference to specific embodiments. It should be understood that the specific embodiments described herein are merely illustrative and not intended to limit the invention. All other embodiments obtained by those skilled in the art based on the embodiments of this invention without inventive effort are within the scope of protection of this invention.

[0021] The main raw materials and reagents used in the following examples and comparative examples have the following sources and specifications. Reagents not specifically mentioned are all commercially available analytical grade or higher grade products.

[0022] The silicon carbide powder is a high-purity α-crystalline monoclinic phase structure, with CAS number 409-21-2, purity greater than or equal to 99.5%, and a D50 median diameter of 1.0 μm to 3.0 μm.

[0023] Niobium pentoxide powder has a monoclinic crystal structure, CAS number 1313-96-8, purity greater than or equal to 99.9%, and a D50 median diameter of 1.0 μm to 5.0 μm.

[0024] Tantalum powder is a high-purity elemental metal powder with CAS number 7440-25-7, a metal purity greater than or equal to 99.9%, and a D50 median particle size distribution of 0.5 μm to 1.5 μm.

[0025] The remaining raw materials, such as cobalt-chromium-molybdenum alloy powder (Co-28Cr-6Mo), high-purity argon, hydrogen-argon mixture, and dental feldspar porcelain powder, are all commercially available medical-grade or industrial-grade products, and their physicochemical properties all meet the corresponding industry standards.

[0026] Preparation Example 1: This preparation example provides a method for preparing a carbon-fixed heterogeneous composite alloy powder, including the following steps: Accurately weigh 100 parts by weight of cobalt-chromium-molybdenum alloy powder (Co-28Cr-6Mo), 0.05 parts by weight of silicon carbide (SiC) powder, 0.15 parts by weight of niobium pentoxide (Nb₂O₅) powder, and 0.25 parts by weight of tantalum (Ta) powder. Place the above four components into the mixing tank of a three-dimensional mixer. After sealing the mixing tank, evacuate it, and then fill it with high-purity argon as a protective atmosphere. Set the mixer speed to 30 rpm and continue dry mixing at room temperature for 4 hours to obtain a carbon-fixed heterogeneous composite alloy powder.

[0027] Preparation Example 2: This preparation example provides a method for preparing a carbon-fixed heterogeneous composite alloy powder, including the following steps: Accurately weigh 100 parts by weight of cobalt-chromium-molybdenum alloy powder (Co-28Cr-6Mo), 0.10 parts by weight of silicon carbide (SiC) powder, 0.28 parts by weight of niobium pentoxide (Nb₂O₅) powder, and 0.48 parts by weight of tantalum (Ta) powder. Place the above four components into the mixing tank of a three-dimensional mixer. After sealing the mixing tank, evacuate it, and then fill it with high-purity argon as a protective atmosphere. Set the mixer speed to 40 rpm and continue dry mixing at room temperature for 5 hours to obtain a carbon-fixed heterogeneous composite alloy powder.

[0028] Preparation Example 3: This preparation example provides a method for preparing a carbon-fixed heterogeneous composite alloy powder, including the following steps: Accurately weigh 100 parts by weight of cobalt-chromium-molybdenum alloy powder (Co-28Cr-6Mo), 0.15 parts by weight of silicon carbide (SiC) powder, 0.40 parts by weight of niobium pentoxide (Nb₂O₅) powder, and 0.70 parts by weight of tantalum (Ta) powder. Place the above four components into the mixing tank of a three-dimensional mixer. After sealing the mixing tank, evacuate it, and then fill it with high-purity argon as a protective atmosphere. Set the mixer speed to 50 rpm and continue dry mixing at room temperature for 6 hours to obtain a carbon-fixed heterogeneous composite alloy powder.

[0029] Example 1: This example provides a method for preparing a porcelain dental restoration. (See attached document.) Figure 1 This includes the following steps: The carbon-fixed heterogeneous composite alloy powder obtained in Preparation Example 1 was added to the powder supply chamber of a selective laser melting (SLM) apparatus. The substrate was preheated to 150°C. The molding chamber was sealed and circulated with high-purity argon gas, maintaining the oxygen volume concentration below 0.1%. The single-layer thickness of the powder spreading device was set to 30 μm. A dual-track scanning was performed based on the slice data: the first track used a focused laser beam with a power of 150 W, a scanning speed of 600 mm / s, a spot diameter of 50 μm, and a scanning interval of 60 μm; immediately following, a second defocused scan was performed along the original path with a laser power of 80 W, a scanning speed of 400 mm / s, and a spot diameter of 150 μm. This process was repeated layer by layer until the metal substrate was formed, followed by powder cleaning.

[0030] The molded substrate was placed in an atmosphere tube furnace equipped with a mass flow controller, and a reducing gas mixture consisting of 10% hydrogen and 90% argon was introduced, with the dew point controlled at -40°C and maintained at atmospheric pressure. The furnace temperature was raised to 800°C at a rate of 5°C / min, held at that temperature for 30 min, and then cooled to room temperature with the furnace.

[0031] Feldspar dental porcelain powder slurry was uniformly applied to the heat-treated substrate surface. The substrate was then transferred to a vacuum porcelain furnace and heated to 550°C at a rate of 30°C / min. Heating was then paused, and a vacuum pump was simultaneously activated to evacuate the substrate to an absolute pressure below 50 Pa. This pressure was maintained isothermally for 10 minutes. Subsequently, this low-pressure state was maintained, and the temperature was increased to 900°C at a rate of 30°C / min, held for 1 minute for melting and sintering, then the vacuum was removed and the substrate was cooled in the furnace.

[0032] Example 2: This example provides a method for preparing a porcelain dental restoration, including the following steps: The carbon-fixed heterogeneous composite alloy powder obtained in Preparation Example 2 was added to the powder supply chamber of a selective laser melting (SLM) apparatus. The substrate was preheated to 175°C. The molding chamber was sealed and circulated with high-purity argon gas, maintaining the oxygen volume concentration below 0.1%. The single-layer thickness of the powder spreading device was set to 40 μm. A dual-track scanning was performed based on the slice data: the first track used a focused laser beam with a power of 200 W, a scanning speed of 800 mm / s, a spot diameter of 65 μm, and a scanning spacing of 75 μm; immediately following, a second defocused scan was performed along the original path with a laser power of 100 W, a scanning speed of 600 mm / s, and a spot diameter of 175 μm. This process was repeated layer by layer until the metal substrate was formed, followed by powder cleaning.

[0033] The molded substrate was placed in an atmosphere tube furnace equipped with a mass flow controller, and a reducing gas mixture consisting of 10% hydrogen and 90% argon by volume was introduced, with the dew point controlled at -35°C and maintained at atmospheric pressure. The furnace temperature was raised to 820°C at a rate of 8°C / min, held at that temperature for 35 minutes, and then cooled to room temperature with the furnace.

[0034] Feldspar dental porcelain powder slurry was uniformly applied to the heat-treated substrate surface. The substrate was then transferred to a vacuum porcelain furnace and heated to 580°C at a rate of 40°C / min. Heating was then paused, and a vacuum pump was simultaneously activated to evacuate the substrate to an absolute pressure below 50 Pa. This pressure was maintained isothermally for 12 minutes. Subsequently, this low-pressure state was maintained, and the temperature was increased to 910°C at a rate of 40°C / min, held for 1.5 minutes for melting and sintering, then the vacuum was removed and the substrate was cooled in the furnace.

[0035] Example 3: This example provides a method for preparing a porcelain dental restoration, including the following steps: The carbon-fixed heterogeneous composite alloy powder obtained in Preparation Example 3 was added to the powder supply chamber of a selective laser melting (SLM) apparatus. The substrate was preheated to 200°C. The molding chamber was sealed and circulated with high-purity argon gas, maintaining the oxygen volume concentration below 0.1%. The single-layer thickness of the powder spreading device was set to 50 μm. A dual-track scanning was performed based on the slice data: the first track used a focused laser beam with a power of 250 W, a scanning speed of 1000 mm / s, a spot diameter of 80 μm, and a scanning interval of 90 μm; immediately following, a second defocused scan was performed along the original path with a laser power of 120 W, a scanning speed of 800 mm / s, and a spot diameter of 200 μm. This process was repeated layer by layer until the metal substrate was formed, followed by powder cleaning.

[0036] The molded substrate was placed in an atmosphere tube furnace equipped with a mass flow controller, and a reducing gas mixture consisting of 10% hydrogen and 90% argon was introduced, with the dew point controlled at -30°C and maintained at atmospheric pressure. The furnace temperature was raised to 850°C at a rate of 10°C / min, held at that temperature for 45 min, and then cooled to room temperature with the furnace.

[0037] Feldspar dental porcelain powder slurry was uniformly applied to the heat-treated substrate surface. The substrate was then transferred to a vacuum porcelain furnace and heated to 600°C at a rate of 50°C / min. Heating was then paused, and a vacuum pump was simultaneously activated to evacuate the substrate to an absolute pressure below 50 Pa. This pressure was maintained isothermally for 15 minutes. Subsequently, this low-pressure state was maintained, and the temperature was increased to 920°C at a rate of 50°C / min, held for 2 minutes for melting and sintering, and then the vacuum was removed and the substrate cooled in the furnace.

[0038] Comparative Example 1: Compared with Example 2, the difference is that no microalloying additives such as silicon carbide, niobium pentoxide and tantalum were added, and only pure cobalt chromium molybdenum alloy powder was used; selective laser melting was performed only once; heat treatment was carried out in a conventional air atmosphere with pre-oxidation at 900°C; the porcelain baking stage was directly and continuously heated, skipping the isothermal degassing stage, and the rest were the same.

[0039] This comparative example serves as a traditional blank control. Due to the lack of the microscopic bonding and modification mechanisms of this invention, its performance is only a baseline. It is used to highlight the synergistic effect and indispensability of the various technical features of this invention in overcoming the pain points of low metal-ceramic bonding strength, high porosity, and interface embrittlement.

[0040] Comparative Example 2: Compared with Example 2, the difference is that the tantalum powder component was removed from the raw materials for preparing the carbon-fixed heterogeneous composite alloy powder, while the rest are the same.

[0041] This comparative example is used to verify the devastating consequences of lacking the tantalum carbon fixation mechanism. The silicon carbide pre-embedded replacement reaction releases a large amount of free carbon. If tantalum is not used to lock it in situ into a stable tantalum carbide dispersed phase, the massive amount of free carbon will directly attack the matrix and bind to chromium in the ultra-high temperature molten pool, resulting in a severe grain boundary chromium depletion effect, which ultimately leads to severe embrittlement of the matrix grain boundaries and a sharp drop in corrosion resistance.

[0042] Comparative Example 3: Compared with Example 2, the difference is that only the first focused laser beam scan is performed in the selected area laser melting stage, and the second defocus scan is removed, while the rest are the same.

[0043] This comparative example is used to verify the consequences of lacking a defocusing rheological slag removal mechanism. The lack of power-reduced, large-spot defocusing scanning makes it impossible to effectively widen the molten pool surface area and prolong the liquid phase residence time without causing matrix boiling. This results in insufficient Marangoni convection uplift time for low-density amorphous silica, leading to the formation of micron-sized brittle inclusions within the metal and causing complete failure of precursor enrichment on the working surface.

[0044] Comparative Example 4: Compared with Example 2, the difference is that the heat treatment atmosphere was changed to a conventional wet reducing mixture with a dew point of about 20°C that was not dehydrated; all other aspects were the same.

[0045] This comparative example is used to verify the consequences of disrupting the thermodynamic selective oxidation mechanism. The extremely high water vapor content in the unremoved gas causes the oxygen partial pressure within the system to far exceed the safe threshold for protecting chromium from oxidation in the Ellingham diagram. This high oxygen partial pressure environment not only oxidizes niobium but also leads to severe uncontrolled oxidation of chromium oxide on the substrate surface, completely destroying the chromium-deficient protective layer and ultimately causing the collapse of the gold-ceramic bond.

[0046] Comparative Example 5: Compared with Example 2, the difference is that the isothermal vacuum holding stage at 580°C was skipped in the porcelain firing stage, and the temperature was directly raised to 910°C at a constant rate for melting and sintering. All other aspects are the same.

[0047] This comparative example is used to verify the consequences of lacking an isothermal kinetic degassing mechanism. The glass transition temperature of feldspar dental porcelain powder is usually between 600°C and 650°C. If Knudsen diffusion degassing is not performed using a high vacuum before this critical point, i.e., at 580°C when the pores are fully interconnected, the porcelain powder will instantly close the pore network once it enters the glassy melting rheological stage, resulting in severe deep gas retention. Ultimately, this leads to vapor explosion and a surge in interfacial porosity at high temperatures.

[0048] Test Example 1: The metal substrate samples of each embodiment and comparative example were removed from the molded substrate using a wire cutting device, and ultrasonically cleaned with anhydrous ethanol for 20 minutes to remove unmelted powder and impurities adhering to the sample surface. Then, they were placed in a vacuum drying oven at 80 degrees Celsius for 2 hours for continuous drying.

[0049] Using an electronic analytical balance with an accuracy of 0.0001 g, the dry mass of each sample in air was weighed and recorded at room temperature (25 degrees Celsius).

[0050] An Archimedes displacement method density testing kit was used, with anhydrous ethanol as the impregnation liquid. The suspended mass of each sample in the liquid was measured and recorded. During the measurement process, it was observed and ensured that there were no attached air bubbles on the sample surface.

[0051] The actual volume and density of each sample were calculated based on the measured mass in air and the suspended mass in the liquid. The relative density and porosity were then extrapolated using the theoretical density value of pure cobalt-chromium-molybdenum alloy. Three parallel samples were randomly selected from each group for repeated testing, and the specific data were compiled.

[0052] Table 1. Test results of density and porosity of metal substrates in each embodiment and comparative example. According to the data in Table 1, Comparative Example 1, which uses pure cobalt-chromium-molybdenum alloy powder for conventional single-pass scanning molding, has an average relative density of 99.13%, demonstrating the basic densification level under conventional processing conditions of this equipment. When silicon carbide and niobium pentoxide are introduced into the powder system but the second defocusing scanning process is eliminated, as shown in Comparative Example 3, the average relative density of the sample significantly decreases to 96.82%, and the porosity increases to 3.18%.

[0053] This indicates that during selective laser melting, a pre-embedded substitution reaction between silicon carbide and niobium pentoxide generates amorphous silicon dioxide. Lacking mechanisms to expand the molten pool surface area and prolong the liquid phase residence time, these low-density silicon dioxide phases cannot effectively overcome the viscous resistance of the liquid metal and fail to float and dissipate before solidification. Ultimately, they remain within the metal matrix as micron-sized inclusions, directly leading to increased internal porosity and decreased macroscopic density.

[0054] In contrast, Examples 1 to 3, which introduced pre-embedded reaction additives and superimposed a second defocusing scanning process, achieved average relative densities of 99.52%, 99.73%, and 99.48%, respectively. These were not only significantly higher than Comparative Example 3, but also better than Comparative Example 1, which was based on a pure alloy baseline.

[0055] This demonstrates that reducing laser power and increasing the spot diameter alters the hydrodynamic state of the molten pool. The lower volume energy density prevents the base metal from boiling while maintaining the surface tension gradient. This state provides sufficient Marangoni convection, allowing the internally generated amorphous silica phase to gain adequate time and driving force for uplift, successfully enriching it on the sample surface. These test results, from a macroscopic physical quantitative perspective, confirm the effectiveness of the dual-channel defocusing rheological slag removal mechanism, indicating that this method can simultaneously ensure and improve the mechanical density of the substrate material during microalloying reactions.

[0056] Test Example 2: The metal substrate samples from each embodiment and comparative example were cut into block samples of uniform size using a metallographic cutting machine. The working surface of the sample was successively polished with silicon carbide wet sandpaper ranging from 400 grit to 2000 grit, followed by ultrasonic cleaning with anhydrous ethanol for ten minutes and drying. One square centimeter of test exposure area was retained, and the remaining non-test areas were sealed with insulating resin.

[0057] Potentiodynamic polarization curves were tested using an electrochemical workstation with a standard three-electrode system. A silver-silver chloride electrode was used as the reference electrode, a platinum sheet as the auxiliary electrode, and the prepared sample as the working electrode. Artificial saliva at 37 degrees Celsius was used as the test solution. After the sample was allowed to stand at its open-circuit potential in the solution for one hour, polarization was measured at a scan rate of 1 millivolt per second, and the corrosion current density was recorded and fitted.

[0058] Static ion release tests were conducted according to dental metal material standards. After polishing and cleaning, samples of the same surface area were immersed in a 37°C test solution containing 0.1% lactic acid and 0.1% sodium chloride, respectively, and then sealed and kept at a constant temperature for seven days.

[0059] The test solution was extracted after soaking for seven days, and the concentrations of cobalt and chromium ions in the solution were determined by inductively coupled plasma mass spectrometry. Based on the exposed surface area of ​​the sample, the total amount of metal ions precipitated per square centimeter of surface area was calculated, and multiple sets of scatter plot data were recorded and the average value was calculated.

[0060] Table 2. Electrochemical corrosion and ion deposition test results of metal substrates in each embodiment and comparative example. According to the data in Table 2, Comparative Example 1, which uses pure cobalt-chromium-molybdenum alloy powder, exhibits basic corrosion resistance, with a self-corrosion current density of 0.68 μA / cm², and cobalt and chromium ion deposition amounts of 5.12 μg / cm² and 2.33 μg / cm², respectively. When silicon carbide and niobium pentoxide are introduced into the powder raw material but the tantalum powder component is removed, as shown in Comparative Example 2, the self-corrosion current density increases sharply to 3.84 μA / cm², and the cobalt and chromium ion deposition amounts surge to 16.75 μg / cm² and 9.41 μg / cm², respectively.

[0061] This indicates that, in the absence of tantalum, the free carbon released from the silicon carbide-niobium pentoxide substitution reaction directly combines with chromium in the matrix within the molten pool, forming a chromium carbide precipitate. This process consumes a large amount of free chromium in the matrix solid solution, creating severely chromium-depleted zones at grain boundaries and phase boundaries. The presence of these chromium-depleted zones disrupts the dense chromium oxide passivation film on the alloy surface, leading to severe intergranular corrosion of the matrix in corrosive media and the dissolution of large amounts of metal ions.

[0062] Comparative Example 4 changed the heat treatment atmosphere to a conventional wet atmosphere without dehydration, and its corrosion current density also increased to 2.75 microamps per square centimeter, and the amount of ion precipitation was significantly higher than the baseline value.

[0063] This proves that the high water content atmosphere causes excessive oxygen partial pressure in the system, which leads to deep and uncontrollable oxidation of chromium on the substrate surface, generating a loose and porous inferior oxide scale, which in turn consumes the chromium reserves inside the substrate and weakens the overall corrosion resistance.

[0064] In contrast, the self-corrosion current density in Examples 1 to 3 remained at a lower level of 0.12 to 0.19 μA per square centimeter, and the amounts of cobalt and chromium ions deposited were significantly lower than those in Comparative Examples 1 and 2. This demonstrates that the tantalum element added to the powder system plays a role in microalloying and carbon fixation during selective laser melting. Because tantalum has a higher affinity for carbon than chromium, tantalum preferentially reacts with free carbon in the system, generating a stable tantalum carbide dispersion phase in situ.

[0065] This mechanism effectively prevents free carbon from plundering chromium, maintaining a uniform distribution of chromium in the cobalt-chromium solid solution. The combination of the carbon fixation mechanism and the ultra-low dew point controlled atmosphere heat treatment allows a dense protective passivation layer to be formed and maintained on the substrate surface.

[0066] Macroscopic test data show that this method, through the in-situ reaction of microalloying elements, not only completes surface modification but also fully guarantees and improves the electrochemical corrosion resistance and clinical ion release safety of the substrate material.

[0067] Test Example 3: Following the standard specifications for the Schwicket crack initiation test of dental porcelain-metal restoration systems, the laser-melted samples from each group were processed into rectangular metal substrates with a length of 25 mm, a width of 3 mm, and a thickness of 0.5 mm. The 8 mm x 3 mm area in the center of the rectangular sample was roughened using alumina blasting material with a particle size of 50 micrometers. The samples were then ultrasonically cleaned in anhydrous ethanol for 10 minutes and dried for later use.

[0068] Within a specific central area after roughening treatment, slurries of opaque porcelain, dentin porcelain, and enamel porcelain were sequentially applied. The coating thickness was strictly controlled using a mold, and the final porcelain layer dimensions after sintering were maintained at 8 mm in length, 3 mm in width, and 1.1 mm in thickness. The specific sintering process for each group of samples strictly followed the temperature and vacuum parameters set in the corresponding examples and comparative examples described above.

[0069] Three-point bending shear bond strength was tested using a microcomputer-controlled electronic universal testing machine. The specimen with the ceramic layer was placed on the test support, with a support span set to 20 mm, the ceramic layer facing downwards and positioned precisely between the two support points. The loading indenter was aligned with the exact center of the metal surface.

[0070] The testing machine was started, and a continuous downward bending load was applied to the metal surface at a constant displacement rate of 1.5 mm per minute until the underlying ceramic layer peeled off or macroscopic fracture occurred. The system automatically recorded the maximum fracture load at the moment of crack initiation. The shear bond strength of the interface was calculated by substituting the values ​​into the Schwicket formula. Four parallel specimens were prepared for each group, and the original test values ​​were recorded and the average value was calculated.

[0071] Table 3. Test results of the metal-ceramic shear bond strength of each embodiment and comparative sample. According to the data in Table 3, the average metal-ceramic shear bond strength of Comparative Example 1, which uses conventional pure cobalt-chromium-molybdenum alloy powder and traditional air pre-oxidation process, is 32.15 MPa. This value is within the normal acceptable level for conventional dental metal materials, and the interfacial adhesion mainly depends on the mechanical and physical bonding brought about by sandblasting and the maintenance of the naturally formed thin oxide film on the substrate surface.

[0072] When the heat treatment atmosphere was changed to a conventional wet reduction mixture without dehydration, as shown in Comparative Example 4, the average metal-ceramic shear bond strength at the interface dropped precipitously to only 17.73 MPa. The data for each parallel sample were generally low and scattered. This indicates that the high dew point in the test environment resulted in a large amount of free moisture in the reaction gas, and the actual oxygen partial pressure in the system significantly exceeded the thermodynamic safety threshold for protecting metallic chromium from oxidation.

[0073] Under these high oxygen partial pressure conditions, chromium on the substrate surface and transition elements leaching from the underlying layer undergo uncontrolled oxidation, forming a thick, loose, porous, and extremely brittle oxide layer. When subjected to the external bending load of a universal testing machine, internal stress cannot be effectively transferred, and the fracture occurs directly within this fragile oxide layer. This typical cohesive failure of the oxide layer leads to the overall spalling of the ceramic block, and the specimen completely loses its physical properties to resist interfacial shear stress.

[0074] The average gold-ceramic shear bond strengths of Examples 1 to 3 reached 48.35 MPa, 52.73 MPa, and 47.73 MPa, respectively, significantly exceeding the traditional technical baseline of Comparative Example 1. This confirms that the extremely low dew point controlled atmosphere set in this scheme successfully implemented the selective oxidation reaction at the thermodynamic level. In an extremely low moisture environment with deep dehydration guarantee, the oxidation tendency of chromium on the substrate surface is limited by thermodynamic conditions, while the active metal niobium enriched to the surface in the early stage, due to its lower oxidation free energy, is transformed in situ into a niobium pentoxide micro-bonded layer.

[0075] During the subsequent high-temperature ceramic firing stage, the active niobium pentoxide layer underwent deep chemical dissolution and ion-level interdiffusion diffusion with the silicon-oxygen tetrahedral glass network in the feldspar ceramic powder, forming a dense chemically bonded transition zone at the metal-ceramic interface.

[0076] Test data confirms that this chemically bonded transition zone completely replaces the fragile thick oxide layer in traditional processes, effectively eliminating micro-stress concentration points and potential crack initiation sites at the interface. Externally applied mechanical shear loads can be smoothly conducted and dispersed from the brittle ceramic shell to the tough metal substrate through this bonding zone, macroscopically manifested as a significant increase in the overall fracture load extremum and shear bond strength of the sample.

[0077] Test Example 4: The metal-based composite ceramic layer samples that have completed the porcelain sintering process were placed in deionized water and ultrasonically cleaned for five minutes to remove surface deposits. The surface moisture was then wiped dry with a lint-free cloth and left to stand at room temperature for 24 hours to allow the residual thermal stress inside the sample to be released naturally.

[0078] Thermal shock fatigue testing of the specimens was conducted using a dual-tank thermal shock test chamber. The temperature of the low-temperature water bath was set to 5 degrees Celsius, and the temperature of the high-temperature water bath was set to 55 degrees Celsius. The single immersion time of the specimens in both the high-temperature and low-temperature baths was set to 30 seconds, and the mechanical transfer time between the two water baths was controlled to be within 2 seconds.

[0079] The equipment is started to continuously cycle through hot and cold. Every 100 cycles, the equipment is paused, and each group of samples is removed and observed under a stereomicroscope with 20x magnification to examine the gold-ceramic interface and the integrity of the ceramic layer surface.

[0080] Record the exact number of cycles at which a through-crack or localized spalling of the bonding interface is first observed on the surface of the ceramic layer for each specimen. If a specimen does not experience structural failure after 5,000 thermal cycles, the testing of that specimen is stopped and recorded as greater than 5,000. Five parallel samples are prepared independently for each group for repeated testing, and the specific number of failure cycles is recorded.

[0081] Table 4. Results of thermal shock fatigue failure cycle tests for each embodiment and comparative sample. According to the data in Table 4, the number of thermal shock fatigue failure cycles of the conventional pure cobalt-chromium-molybdenum alloy and the conventional process in Comparative Example 1 ranged from 1680 to 2120, representing the basic thermal fatigue resistance level of the material system under conventional preparation conditions.

[0082] In Comparative Example 5, the powder formulation and heat treatment process were the same as in Example 2, but the isothermal vacuum holding process at 580 degrees Celsius was omitted during the porcelain firing stage, and the temperature was directly raised to 910 degrees Celsius at a constant rate. The test results showed extremely poor thermal shock resistance; all five parallel samples cracked or the porcelain layer peeled off within an extremely low cycle range of 280 to 520 cycles, and the data dispersion was high.

[0083] This indicates that the glass transition temperature of feldspar dental porcelain powders is typically between 600 and 650 degrees Celsius. Without isothermal degassing, continuous heating causes the porous solid porcelain powder to rapidly cross its glass transition point, transforming into a fluid molten glass phase. This rapid transition results in the liquid phase instantly sealing off the previously interconnected internal pore network, blocking the outward diffusion channels of residual gas at the underlying material and interface.

[0084] The trapped gas expands in volume as the temperature rises to 910 degrees Celsius, forming tiny closed-pore defects at the solidified metal-ceramic interface. During frequent alternating heating and cooling tests, these interfacial pores become the initiation points for thermal stress concentration and cracking. The alternating shear stress caused by the difference in thermal expansion coefficients of the different materials is superimposed at the pores, directly accelerating the fatigue failure of the interface.

[0085] The number of thermal shock fatigue failure cycles in Examples 1 to 3 was significantly increased, generally reaching around 4000 cycles, with some samples in Example 2 even withstanding more than 5000 thermal shock cycles. This demonstrates that the operation of setting an isothermal constant pressure at 580 degrees Celsius and simultaneously drawing a vacuum accurately utilized the microporous state of the ceramic powder before melting.

[0086] Within this specific temperature range, the powder particles have not yet undergone viscous flow, and the gaps between the particles constitute open venting channels. Applying an absolute low-pressure environment below 50 Pa allows the gases at the interface and in the underlying layers to be completely extracted from the restoration system via the Knudsen diffusion mechanism. After the gases are exhausted, continued heating and subsequent vitrification melting sintering form a completely dense and pore-free interface structure. This highly dense interface eliminates structural weaknesses and effectively conducts and disperses the temperature gradient stress induced by alternating heating and cooling, demonstrating, from a kinetic heat treatment perspective, the crucial role of isothermal degassing in improving the reliability of restorations during service.

[0087] Test Example 5: After selective laser melting and heat treatment, the metal samples from each group were machined into standard dumbbell-shaped room temperature tensile specimens using wire cutting. The gauge length of the specimen was set to 25 mm, the width of the rectangular cross-section was 5 mm, and the thickness was 1.5 mm. The surface of the specimen was polished step by step with silicon carbide wet sandpaper until visible machining marks were eliminated, and then ultrasonically cleaned in anhydrous ethanol for 10 minutes and dried for later use.

[0088] The tensile specimen is vertically clamped in the upper and lower wedge-shaped grips of the computer-controlled electronic universal testing machine, ensuring that the longitudinal axis of the specimen is strictly aligned with the loading axis. A static extensometer with a gauge length of 20 mm is installed in the gauge length section of the specimen to continuously acquire real strain data during the tensile process.

[0089] The tensile loading rate of the testing machine was set to one millimeter per minute. At room temperature of 25 degrees Celsius, a unidirectional continuous tensile load was applied along the longitudinal axis of the specimen, and the load and displacement data were recorded throughout the process until the specimen underwent macroscopic fracture.

[0090] Extract the stress-strain curve generated by the testing system, calculate the specified non-proportional elongation strength at which the specimen exhibits 0.2% residual deformation, and record it as the yield strength. After fracture, splice the specimens together, measure the final length of the gauge length, calculate the percentage of elongation relative to the original gauge length, and record it as the fracture elongation. Prepare four parallel specimens independently for each group and repeat the test, recording the original data and calculating the average value.

[0091] Table 5. Results of room temperature tensile mechanical properties of metal substrates in each embodiment and comparative example. According to the data in Table 5, Comparative Example 1, which uses pure cobalt-chromium-molybdenum alloy powder for conventional printing, has an average yield strength of 549.0 MPa and an average elongation at break of 8.60%. This data reflects the inherent mechanical characteristics of the cobalt-chromium-molybdenum alloy without added modifiers under conventional selective laser melting (SLM) processes.

[0092] When tantalum powder was removed from the powdered raw material, as shown in Comparative Example 2, the average yield strength of the sample increased slightly to 607.2 MPa due to the precipitation of the carbide phase, but its average elongation at break dropped sharply to 2.15%, and the sample underwent brittle fracture after minimal plastic deformation. This indicates that after silicon carbide and niobium pentoxide undergo a pre-embedded replacement reaction in the molten pool, a large amount of highly active free carbon is released. In the absence of tantalum as a carbon scavenger, this free carbon directly combines with chromium in the matrix solid solution, resulting in the precipitation of a large amount of network-distributed hard and brittle chromium carbide phase at the grain boundaries.

[0093] The continuous enrichment of such grain boundary carbides severely hinders the slip movement of dislocations, triggering a typical grain boundary embrittlement effect. This causes the material to directly initiate cracks along the grain boundaries and propagate rapidly when subjected to macroscopic tension, resulting in the loss of its original ductility.

[0094] Comparative Example 3, while maintaining the addition of all components, eliminated the second defocusing scanning process, resulting in a decrease in average yield strength to 501.5 MPa and elongation at break to 4.35%. This demonstrates that without the large-area stable Marangoni convection provided by the defocusing scan, the low-density amorphous silica phase generated by the substitution reaction cannot float to the surface and be discharged in time. A large amount of silica remains inside the metal as micron-sized brittle inclusions, disrupting the continuity of the matrix structure.

[0095] Under external tensile load, the deformation of these internal inclusions at the interface with the metal matrix is ​​inconsistent, becoming dense stress concentration points, which promotes the early initiation of microcracks and causes the specimen to fracture prematurely under low stress.

[0096] The mechanical properties of Examples 1 to 3 are significantly superior to those of the comparative examples described above. Taking Example 2 as an example, its average yield strength reached 737.7 MPa, while the average elongation at break increased to 13.92%, achieving simultaneous growth in strength and plasticity. This confirms that the introduction of tantalum can preferentially and stably combine with the free carbon generated in the reaction, forming a fine and dispersed tantalum carbide solid carbon phase in situ within the grain. This process not only completely avoids the consumption of chromium and grain boundary embrittlement, but also hinders dislocation movement through a second-phase dispersion strengthening mechanism, thereby improving the yield strength.

[0097] Meanwhile, the dual-channel defocusing scanning process effectively eliminated internal silica inclusions, ensuring the high density and structural continuity of the substrate. Test data, from a macroscopic mechanical perspective, verified the synergistic mechanism and practical effectiveness of in-situ solid carbon microalloying and rheological slag removal processes in improving the structural strength and toughness of metal substrates.

[0098] Although embodiments of the invention have been shown and described, it will be understood by those skilled in the art that various changes, modifications, substitutions and alterations can be made to these embodiments without departing from the principles and spirit of the invention, the scope of which is defined by the appended claims and their equivalents.

Claims

1. A cobalt-chromium alloy laser-clad porcelain denture base, characterized in that, It contains the following ingredients by weight: 100 parts of cobalt-chromium-molybdenum alloy powder; 0.05-0.15 parts of silicon carbide; 0.15-0.40 parts of niobium pentoxide; 0.35-0.65 parts of tantalum powder; The substrate contains a tantalum carbide phase that is used to lock free carbon in situ within its internal microstructure, and the surface of the substrate has a niobium pentoxide transition layer that blocks the oxidation of chromium and forms chemical bonds with the ceramic powder.

2. The cobalt-chromium alloy laser-clad porcelain denture base according to claim 1, characterized in that: The cobalt-chromium-molybdenum alloy powder, silicon carbide, niobium pentoxide, and tantalum powder all have spherical morphology, and the particle size distribution range of the powder is 15-53 μm.

3. The cobalt-chromium alloy laser-clad porcelain denture base according to claim 1, characterized in that: The substrate has no micron-sized amorphous silica inclusions in its internal metal matrix, and the macroscopic relative density of the substrate is greater than or equal to 99.40%.

4. The cobalt-chromium alloy laser-clad porcelain denture base according to claim 1, characterized in that: The substrate has a room temperature yield strength greater than or equal to 690 MPa and an elongation at break greater than or equal to 11.0%.

5. The cobalt-chromium alloy laser-clad porcelain denture base according to claim 1, characterized in that: The surface of the substrate is also coated with a feldspar dental porcelain layer, wherein the silicon-oxygen tetrahedral glass network at the bottom of the feldspar dental porcelain layer and the niobium pentoxide transition layer on the surface of the substrate have an ion-level interdiffusion diffusion structure.

6. A method for preparing a cobalt-chromium alloy laser-clad porcelain denture substrate, comprising the cobalt-chromium alloy laser-clad porcelain denture substrate according to any one of claims 1-5, characterized in that, Includes the following steps: S1. Powder preparation: Weigh out cobalt-chromium-molybdenum alloy powder, silicon carbide, niobium pentoxide and tantalum powder according to the weight ratio, and mix them evenly to obtain composite alloy powder; S2. Selective laser melting and forming: The composite alloy powder is introduced into a selective laser melting device, and the first focused scanning melting is performed on the single-layer spread powder. Then, the second defocusing scanning is performed in situ to induce Marangoni convection for rheological slag removal. The metal substrate is stacked layer by layer to obtain the initial product. S3. Controlled atmosphere heat treatment: The metal substrate sample is placed in a reducing mixed gas atmosphere with an extremely low dew point for heat treatment, which triggers thermodynamic selective oxidation and generates a niobium pentoxide transition layer in situ on the substrate surface to obtain a modified metal substrate. S4. Isothermal degassing porcelain: Dental porcelain powder slurry is applied to the surface of the modified metal substrate. During the heating and sintering process, isothermal vacuum degassing is performed before the porcelain powder crosses the glass transition temperature. Then, the temperature is continued to rise to complete the melting and sintering, thus obtaining the final product.

7. The method for preparing a cobalt-chromium alloy laser cladding porcelain denture substrate according to claim 6, characterized in that: In step S1, silicon carbide, niobium pentoxide, and tantalum powder are pre-treated as follows: the weighed silicon carbide, niobium pentoxide, and tantalum powder are placed together in a planetary ball mill, and anhydrous ethanol is used as the grinding medium. They are pre-mixed and dispersed at a speed of 250-350 rpm for 2-4 hours. After being vacuum dried at 80℃, they are added to a three-dimensional mixer with cobalt-chromium-molybdenum alloy powder for dry mixing.

8. The method for preparing a cobalt-chromium alloy laser cladding porcelain denture substrate according to claim 6, characterized in that: The specific implementation method of step S2 is as follows: The process parameters for the first focusing scan are: laser power 180-220W, scanning speed 800-1000mm / s, and spot diameter 50-80μm; The process parameters for the second defocus scanning are as follows: the laser power is reduced to 80-120W, the scanning speed is reduced to 400-600mm / s, and the spot diameter is expanded to 120-180μm through positive defocus.

9. The method for preparing a cobalt-chromium alloy laser cladding porcelain denture substrate according to claim 6, characterized in that: In step S3, the reducing mixed atmosphere is composed of argon and hydrogen with a volume fraction of 5-10%, and the dew point of the reducing mixed atmosphere is strictly controlled to be lower than or equal to -40℃; the heating rate of the heat treatment is controlled to be 10-15℃ / min, the holding temperature is 900-1000℃, and the holding time is 1-2h.

10. The method for preparing a cobalt-chromium alloy laser cladding porcelain denture substrate according to claim 6, characterized in that: In step S4, the specific implementation method of the isothermal vacuum degassing process is as follows: When the ceramic furnace reaches 580℃, stop heating and maintain an isothermal state at 580℃ while simultaneously drawing a vacuum from the furnace cavity until the absolute pressure in the furnace cavity is below 50Pa. The pressure holding and degassing time should be controlled at 10-15 minutes.

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