【作者】 赵丹；

【导师】 张长瑞； 胡海峰；

【作者基本信息】 国防科学技术大学 ， 材料科学与工程， 2011， 博士

【摘要】 高超音速飞行器速度的提高对飞行器的鼻锥、机翼前缘和发动机等热结构部件的承载和超高温抗烧蚀性能提出了严峻的挑战,因此迫切需要开发具有良好力学性能和抗烧蚀性能的耐超高温材料。针对现有耐超高温材料的不足,提出采用先驱体浸渍-裂解工艺（PIP）制备连续纤维增强耐超高温陶瓷基复合材料。采用混合-反应法制备了TiC、ZrC、ZrB₂等耐超高温陶瓷先驱体,在此基础上采用PIP工艺制备了C/ZrC、C/ZrC-SiC复合材料,研究了界面涂层对复合材料性能的影响,探讨了复合材料在氧乙炔焰和电弧风洞环境中的烧蚀机理。采用混合-反应法制备了TiC、ZrC、ZrB₂等耐超高温陶瓷先驱体,并研究了交联、裂解机理。首先,以钛酸丁酯和二乙烯基苯（DVB）分别为钛源和碳源,混合得到了TiC先驱体。该先驱体在120²10℃可通过DVB自身的交联而固化,283²97℃分解得到t-TiO₂,445⁴61℃继续分解得到裂解碳,1200℃通过碳热还原反应得到TiC陶瓷。提高先驱体中DVB的含量或提高裂解温度,可以提高裂解产物中TiC的结晶度。其次,以锆酸丁酯和DVB分别为锆源和碳源,混合得到了醇基ZrC先驱体。该先驱体在120²10℃通过DVB自身的交联而固化,交联温度达150℃时先驱体部分分解得到无定形ZrO₂,升温裂解时先驱体继续分解得到无定形ZrO₂,400℃时无定形ZrO₂结晶得到t-ZrO₂,400⁵00℃时DVB分解成裂解碳,1400℃以上发生碳热还原反应得到ZrC。随着先驱体中DVB含量的增加和1600℃保温时间的延长,产物中ZrC的结晶度提高。当先驱体中Zr/C比为1/3时,在1600℃裂解2h得到不含氧化物的ZrC陶瓷。第三,以ZrOCl2·8H2O和蔗糖分别为锆源和碳源,制备了水基ZrC先驱体。该先驱体在70℃浓缩干燥,400℃裂解时得到ZrO₂和裂解碳的混合物,1600℃发生碳热还原反应生成ZrC陶瓷。第四,以ZrOCl2·8H2O、硼酸及酚醛树脂为原料,制备了ZrB₂陶瓷先驱体。先驱体溶液在40℃干燥时先转化为溶胶,然后凝胶,最后变成疏松的固体。交联后的先驱体在400℃以下无机化得到ZrO₂、B₂O3和碳,在1600℃陶瓷化得到ZrB₂陶瓷。以上四种先驱体中,TiC先驱体和醇基ZrC先驱体具有较好的交联、裂解特性,能够满足PIP工艺使用要求。以Zr（OBu）4和DVB的混合物为先驱体,采用PIP工艺制备了C/ZrC复合材料。为了提高复合材料的致密化效率和材料性能,确定了低温无机化、高温热处理进行碳热还原的工艺。优化的工艺参数如下:150℃交联、700℃无机化、1600℃热处理、20个PIP致密化周期,制备的C/ZrC复合材料弯曲强度253.6MPa,弹性模量42.3GPa,断裂韧性14.54MPa·m1/2。氧乙炔焰烧蚀300s后,质量烧蚀率为0.0059g/s,线烧蚀率0.0040g/s。采用CVI和PIP工艺在碳纤维表面制备了CVI-SiC、PIP-SiC和PIP-C界面涂层,考察了界面涂层对C/ZrC复合材料性能的影响。研究表明,CVI-SiC涂层后C/ZrC复合材料的力学性能下降,PIP-SiC和PIP-C涂层后C/ZrC复合材料的力学性能提高。引入CVI-SiC和PIP-SiC涂层后,改善了复合材料的抗烧蚀性能。其中,以制备了两个周期PIP-SiC界面层的SiCPIP2-C/ZrC复合材料综合性能最优,其弯曲强度、弯曲模量和断裂韧性分别为319.2MPa、46.3GPa、18.81MPa·m1/2,氧乙炔焰质量烧蚀率和线烧蚀率分别为0.0098g/s和0.0089mm/s。为了提高耐超高温复合材料的抗氧化性能,制备了C/ZrC-SiC复相陶瓷基复合材料。1200℃氧化0.5h后C/ZrC-SiC复合材料的强度和模量分别达到157.5MPa和22.0GPa,静态抗氧化性能远高于C/ZrC复合材料。C/ZrC-SiC复合材料的强度为322.0MPa,模量48.3GPa,断裂韧性为11.55 MPa·m1/2,氧乙炔焰质量烧蚀率和线烧蚀率分别为0.0089g/s和0.0136mm/s。电弧风洞质量烧蚀率和线烧蚀率分别为0.0181g/s和0.0037mm/s。由于SiC氧化形成的低粘度SiO₂熔体在复合材料表面形成了连续的SiO₂保护膜,所以C/ZrC-SiC复合材料在氧乙炔焰环境下的质量烧蚀率比SiCPIP2-C/ZrC复合材料有所下降,线烧蚀率升高,但电弧风洞环境下的抗烧蚀性能下降。研究了C/ZrC复合材料在氧乙炔焰环境下的烧蚀机理。在氧乙炔焰烧蚀考核中,C/ZrC复合材料主要发生了热化学烧蚀,烧蚀过程中发生的主要反应是ZrC及C的氧化,生成的主要产物是ZrO₂和CO₂。烧蚀后C/ZrC复合材料沿厚度方向主要分为四层:表面熔融层、珊瑚状ZrO₂层、氧化层及未烧蚀层,表面熔融层又可分为熔融区、疏松区以及边缘区。表面熔融层各区域的微观形貌虽然存在差异,但是都有ZrO₂熔融,熔融的ZrO₂密封了烧蚀层中的部分孔隙,为内层的C/ZrC复合材料提供抗烧蚀保护。根据烧蚀过程中的温度变化,C/ZrC复合材料的烧蚀主要分为两个阶段,第一阶段温度较低,主要为ZrC和碳纤维的氧化,烧蚀由ZrC和C的氧化反应速率控制。第二阶段温度较高,ZrO₂熔融并铺展,烧蚀速率主要由氧在ZrO₂熔融层中的扩散速率控制。研究了PIP-SiC界面涂层的C/ZrC-SiC复合材料氧乙炔焰烧蚀机理。C/ZrC-SiC复合材料在氧乙炔焰环境下的烧蚀主要是热化学烧蚀和气流冲刷,发生的主要反应是C、ZrC、SiC的氧化,生成的主要产物是CO₂、ZrO₂、SiO₂。随着离烧蚀中心的距离增加,复合材料的烧蚀面依次为中心熔融区、SiO₂耗尽区、SiO₂富集区。沿厚度方向,C/ZrC-SiC复合材料可分为熔融泥浆层、颗粒状SiO₂耗尽层、熔融SiO₂富集层和未烧蚀复合材料层。其中SiO₂富集层连续而致密,为C/ZrC-SiC复合材料提供抗烧蚀保护。随着烧蚀温度的逐渐上升,C/ZrC-SiC复合材料烧蚀过程主要分为三个阶段:首先是较低温度下C、ZrC、SiC的氧化,形成疏松的氧化层,该阶段的烧蚀速率由C、ZrC、SiC与氧之间的反应速率控制;随后,温度达到SiO₂的熔点,SiO₂融化并在气流冲刷下铺展,在复合材料表面形成保护层,同时低粘度的SiO₂熔体在高速气流的冲刷下流失,该阶段的烧蚀由氧通过SiO₂熔融层的扩散速率以及SiO₂的冲刷流失速率控制;当温度达到SiO₂的沸点以后,SiO₂开始从表面挥发,该阶段的烧蚀由氧通过SiO₂熔融层的扩散速率、SiO₂的冲刷流失速率和SiO₂的挥发速率控制。研究了电弧风洞考核环境下耐超高温陶瓷基复合材料的烧蚀机理。C/ZrC复合材料在电弧风洞烧蚀过程中形成了连续而致密的ZrO₂熔融层,并粘附于烧蚀表面,保护复合材料,材料具有较好的抗烧蚀性能,其烧蚀主要是化学烧蚀。C/ZrC-SiC复合材料在电弧烧蚀过程中生成了粘度较低的SiO₂-ZrO₂混合物,从烧蚀表面流失。熔融层的消耗导致内部材料暴露在烧蚀气流中,复合材料被严重烧蚀。C/ZrC-SiC复合材料在电弧风洞考核中的烧蚀主要是化学烧蚀和机械冲刷。更多还原

【Abstract】 Ultra high temperature materials with excellent mechanical and ablation properties were needed for the nose caps, sharp leading edges and rocket engines of hypersonic aerospace vehicles. The proposal to prepare continuous fiber reinforced ultra high temperature ceramics matrix composites by precursor infiltration and pyrolysis （PIP） process is presented to improve the properties of the existing ultra high temperature materials. TiC, ZrC and ZrB₂ precursors were firstly prepared by mixing-reaction method, then by use of these precursors C/ZrC and C/ZrC-SiC composites were fabricated, accordingly the influence of interface coatings on the properties of the composites, and the ablation mechanisms of the composites in arc-jet wind tunnel and oxyacetylene flame environments were investigated and discussed.TiC, ZrC and ZrB₂ precursors were prepared by mixing-reaction method, and the cross-linking and pyrolysis mechanisms were investigated. Firstly, TiC precursor was obtained by mixing titanium butoxide and divinylbenzene （DVB）. The precursor cross-linked at 120²10oC via the cross-linking of DVB, and decomposed into t-TiO₂ and carbon at 283²97oC and 445⁴61oC, respectively. TiC was formed at 1200oC by carbo-thermal reaction. The crystalline degree of TiC increased while the content of DVB increased and/or pyrolysis temperature increased. Secondly, hydroxyl-ZrC precursor was prepared with zirconium butoxide and DVB as sources of zirconium and carbon, respectively. The precursor cross-linked at 120²10oC, and partialy decomposed into amorphous ZrO₂ at above 150oC. Amorphous ZrO₂ changed into t-ZrO₂ at 400oC. DVB decomposed into carbon at 400⁵00oC. ZrC was obtained by carbo-thermal reaction at temperature higher than 1400oC. The crystalline degree of ZrC increased while the content of DVB increased or the holding time at 1600oC increased. ZrC without oxide was obtained when the precursor with a Zr/C ratio of 1/3 was pyrolyzed at 1600oC for 2 hours. Thirdly, aqueous-ZrC precursor was prepared with zirconium oxychloride （ZrOCl2·8H2O） and sucrose as sources of zirconium and carbon, respectively. This precursor was dried at 70oC, and pyrolyzed at 400oC to obtain the blends of ZrO₂ and carbon, and at 1600oC to obtain ZrC by carbo-thermal reaction. Fourthly, ZrB₂ precursor was prepared with ZrOCl2·8H2O, boric acid and phenolic resin as sources of zirconium, boron and carbon, respectively. When dried at 40oC, the precursor turned into sol, then gel, and at last loose solid. The precursor decomposed into ZrO₂, B₂O3 and carbon under 400oC, and turned into ZrB₂ at 1600oC. The TiC precursor and hydroxyl-ZrC precursor had better cross-linking and pyrolysis properties, and met the requirements of PIP process.C/ZrC composite was fabricated by PIP process while the mixture of Zr（OBu）4 and DVB was used as ZrC precursor. A process of inorganic treatment at low temperature and carbo-thermal treatment at high temperature was established to improve the densification efficiency and the properties of the composite. The parameters were optimized as cross-linking at 150oC, inorganic treatment at 700oC, carbo-thermal treatment at 1600oC, and 20 PIP cycles. The as-obtained C/ZrC composite had a tensile strength of 253.6 MPa, elastic modulus of 42.3 GPa, and fracture toughness of 14.54 MPa·m1/2. After ablated by oxyacetylene flame for 300 seconds, the C/ZrC composite showed a mass loss rate of 0.0059 g/s and a linear recession rate of 0.0040 mm/s.CVI-SiC, PIP-SiC, PIP-C interface coatings were fabricated on carbon fiber by CVI and PIP processes, respectively. The mechanical properties of the C/ZrC composites with CVI-SiC interface coatings decreased, and those of the C/ZrC composites with PIP-SiC and PIP-C interface coatings increased. Furthermore, the CVI-SiC and PIP-SiC interface coatings improved the ablation properties of the C/ZrC composite. The SiCPIP2-C/ZrC composite with PIP-SiC interface coating of 2 cycles had the best properties: a tensile strength of 319.2 MPa, elastic modulus of 46.3 GPa, fracture toughness of 18.81 MPa·m1/2, mass loss rate of 0.0098 g/s and linear recession rate of 0.0089 mm/s during the oxyacetylene torch test.C/ZrC-SiC composite was fabricated to improve the oxidation property of the composite. After oxidized at 1200oC for 0.5 hour, the C/ZrC-SiC composite showed a tensile strength of 157.5 MPa and an elastic modulus of 22.0 GPa, which were much better than those of C/ZrC composite. The tensile strength of the C/ZrC-SiC composite was 322.0 MPa, the elastic modulus was 48.3 GPa, the fracture toughness was 11.55 MPa·m1/2. The mass loss rate and linear recession rate of the C/ZrC-SiC composite were 0.0089 g/s and 0.0136 mm/s in oxyacetylene flame environment, respectively, while these were 0.0181 g/s and 0.0037 mm/s in arc-jet wind tunnel environment, respectively. The SiO₂ melt layer on the ablation surface from the SiC oxidation increases the mass loss rate but decreases the linear recession rate of the C/ZrC-SiC composite. When ablated in arc-jet wind tunnel environment, the C/ZrC-SiC composite showed a worse ablation property than C/ZrC composite.The ablation mechanism of the C/ZrC composite in oxyacetylene flame environment was discussed. Thermal-chemical process was the main ablation mechanism during the oxyacetylene torch test. ZrC and carbon was oxidized into ZrO₂ and CO₂. After the ablation, the composite was divided into four layers as the melting layer, the loose tree-coral-like ZrO₂ layer, the undersurface oxidation layer and the composite layer. The melting layer was divided into melting, porous and marginal region. Though slight difficulties existed in the different regions, ZrO₂ melt existed in all regions and partially sealed the pores and protected the composite. The ablation process may be divided into two steps according to the temperature gradient. Firstly, oxidation of carbon fiber and ZrC matrix began at relative lower temperatures to form a porous structure. In this step the ablation rate was controlled by the oxidation rate of carbon and ZrC. Secondly, ZrO₂ melted and spread on the ablation surface while the temperature was raised. The ablation rate was controlled by the diffusion rate of oxygen through the ZrO₂ melting layer.The ablation mechanism of C/ZrC-SiC composite with PIP-SiC interface coating in oxyacetylene flame environment was investigated. The ablation of C/ZrC-SiC composite was a combination of thermal-chemical ablation and mechanical erosion. ZrC, SiC and carbon were oxidized into ZrO₂, SiO₂ and CO₂ respectively. The ablated surface was divided into a melting region, a SiO₂ depleted region and a SiO₂ enriched region. Through thickness the ablated composite was divided into a slurry layer, a grain-like SiO₂ depleted layer, a melting SiO₂ enriched layer and a composite layer. The melting SiO₂ enriched layer was dense and integral, which protected the composite. With temperature increase the ablation process of C/ZrC-SiC composite may be divided into three steps. Firstly, oxidation of carbon fiber, SiC and ZrC matrix began at relative lower temperature to form a porous structure. In this step the ablation rate was controlled by the oxidation rate. Secondly, while the temperature was raised, SiO₂ melted and spread on the ablated surface and was blown away by the gas flux. The ablation rate was controlled by the diffusion rate of oxygen through the SiO₂ melting layer and the eroding rate. Thirdly, SiO₂ evaporated after the temperature reached the boiling point of SiO₂. The ablation rate was controlled by the diffusion rate of oxygen, the evaporate rate of SiO₂ and the eroding rate.The ablation mechanisms of the ultra high temperature composites during arc-jet wind tunnel test were also investigated. For C/ZrC composite, dense ZrO₂ melting layer formed during the arc-jet wind tunnel test, adhered to the ablation surface firmly, and protected the composite. Thus, the C/ZrC composites showed an excellent ablation property, and chemical reaction was the main mechanism during the arc-jet wind tunnel test. When C/ZrC-SiC composite was ablated by arc-jet wind tunnel, low viscosity ZrO₂-SiO₂ melting layer formed and was blown away from the surface. The composite was exposed to the ablation flow and ablated seriously. Chemical reaction and mechanical erosion was the main mechanism of the C/ZrC-SiC composite during the arc-jet wind tunnel test.更多还原