碳纤维增韧陶瓷基复合材料高温氧化性能研究进展：氧化机制、氧化损伤实验与模型

方国东; 王章文; 李赛; 王兵; 孟松鹤

doi:10.13801/j.cnki.fhclxb.20240418.004

碳纤维增韧陶瓷基复合材料高温氧化性能研究进展：氧化机制、氧化损伤实验与模型

哈尔滨工业大学　复合材料与结构研究所，哈尔滨 150001

基金项目: 国家自然科学基金(12090034)；黑龙江省自然科学基金(YQ2021A004)

详细信息

通讯作者:
方国东，博士，教授，博士生导师，研究方向为编织复合材料力学和多场耦合数值计算　E-mail: fanggd@hit.edu.cn

中图分类号: TB332
计量
- 文章访问数: 838
- HTML全文浏览量: 526
- PDF下载量: 161
出版历程
- 收稿日期: 2024-02-21
- 修回日期: 2024-03-19
- 录用日期: 2024-04-02
- 网络出版日期: 2024-04-18
- 刊出日期: 2024-08-31

Review of high-temperature oxidation properties for carbon fiber toughened ceramic matrix composites: Oxidation mechanisms, oxidation damage experiments and models

Center for Composite Materials and Structure, Harbin Institute of Technology, Harbin 150001, China

Funds: National Natural Science Foundation of China (12090034); Natural Science Foundation of Heilongjiang Province (YQ2021A004)

HTML全文

摘要

摘要: 碳纤维增韧陶瓷基复合材料兼具陶瓷材料优良的抗氧化腐蚀性能和碳纤维材料增强增韧的力学性能，已成为最有潜力的高超声速飞行器热防护候选材料。碳纤维增韧陶瓷基复合材料在多物理场耦合服役环境下的高温氧化损伤失效机制对热防护材料设计及性能表征与评价至关重要，也一直是国内外学者研究的热点。本文从高温氧化机制、耦合失效实验、高温氧化模型3个方面对C/SiC和C/ZrB₂-SiC复合材料进行详细论述和总结，对相应的研究方法的局限性和适用范围进行分析和评价，并展望了碳纤维增韧陶瓷基复合材料氧化损伤研究的发展趋势，进而为碳纤维增韧陶瓷基复合材料在热力氧耦合条件下的热/力响应及性能评价研究起到指导作用。
- 碳纤维增韧陶瓷基复合材料 /
- 氧化机制 /
- 损伤失效 /
- 实验 /
- 氧化模型
Abstract: Carbon fiber toughened ceramic matrix composites inherit the excellent mechanical properties of carbon fibers and the high oxidation and corrosion resistance of ceramics, become the most promising candidate thermal protection materials for hypersonic vehicles. The high-temperature oxidation mechanisms and damage mechanical behaviors of carbon fiber toughened ceramic matrix composites in coupled service environments are important topics in the study of design, property characteristic and evaluation of thermal protection material. This provides a detailed discussion and summary of the research and analysis methods employed to characterize the oxidation damage of C/SiC and C/ZrB₂-SiC composites in three aspects: High-temperature oxidation mechanisms, coupled failure experiment, and high-temperature oxidation model. The limitations and applicability of various research methods are analyzed and evaluated. In addition, the development trend of investigation of oxidation damage in carbon fiber toughened ceramic matrix composites is provided. It provides guidance for the study of thermal/mechanical response analysis and performance evaluation of carbon fiber toughened ceramic matrix composites in the thermal/mechanical/oxygen environment.
- carbon fiber toughened ceramic matrix composites /
- oxidation mechanism /
- damage failure /
- experiment /
- oxidation model

HTML全文

可再生能源的间隙性和波动性为能量的安全、高效利用带来较大挑战，已经成为经济可持续发展面临的重要议题^[1]。钒液流电池(Vanadium liquid flow battery，VRB)可将电能高效储存，具有本征安全、响应速度快和运行寿命长等特点，相对于其他类型储能技术，钒电池技术在大规模、长时储能领域显示出优越的发展前景^[1-3]。

作为钒电池的关键部件，质子交换膜(Proton exchange membrane，PEM)需具有如下特性^[4-5]：(1)高离子选择性，即高质子传导率和低钒离子渗透率；(2)高拉伸强度和穿刺强度；(3)良好的理化稳定性；(4)较好的经济性。目前，基于全氟磺酸树脂的PEM在VRB中的应用最为广泛，如杜邦公司Nafion系列膜，但其溶胀性和离子选择性等与VRB的严格要求仍存在差距，优化膜的离子渗透性是本领域的研究热点^[6-8]。实验表明，向膜中引入具有适宜孔道结构和表面性质的功能材料是改善其离子选择性的有效手段^[5]，如金属-有机框架(Metal-organic frameworks，MOFs)材料^[9-11]。

在MOFs材料中，UIO-66 (Zr-MOF)由六核氧化锆簇作为二级构建单元和1, 4-苯二甲酸接头构建^[12]，具有刚性结构和较强的耐酸性，同时拥有介于水和离子(＜0.3 nm)和钒离子(＞0.6 nm)的孔道尺寸^[13]。UIO-66主要通过水热合成法制备，但在合成过程中，烘箱加热较难提供均匀的受热环境，存在合成时间长、材料结构均一性差等问题^[14]。微波加热可促进反应物中分子或离子直接耦合，实现能量的快速传导，具有反应速率快和产率高等优点，已被广泛用于多种纳米材料的合成^[14-16]。Zhai等^[6]通过在磺化聚醚醚酮(SPEEK)中掺入15wt% 的传统水热法UIO-66-NH₂来改善复合膜的质子传导率，在120 mA·cm⁻²电流密度下，电池的能量效率(Energy efficiency，EE)为77.3%。Lu等^[17]制备高质子选择性的聚多巴胺(PDA)@MOF-808，并将其掺入SPEEK中以优化复合膜的性能，在120 mA·cm⁻²电流密度下，该膜的EE为83.9%。贾儒等^[13]和Wan等^[14]也分别对UIO-66的合成和应用展开研究，均采用耗时较长的传统水热法。

鉴于此，本文采用不同加热方式制备氨基官能化的UIO-66 (UIO-66-NH₂)，验证微波加热在UIO-66-NH₂材料的合成效率、结构优化方面的优势，并将其与Nafion 掺杂，通过溶液浇筑法制备复合膜，对膜的理化性质和电池性能进行表征，探讨新型膜材料对钒电池性能的影响。

1. 实验部分

1.1 原材料

氯化锆(ZrCl₄，98%)、2-氨基对苯二甲酸(BDC-NH₂，99%)，上海麦克林生化科技有限公司；Nafion溶液(固含量5%)，美国杜邦公司；无水乙醇，沈阳试剂二厂；N, N-二甲基甲酰胺(DMF)、MgSO₄、硫酸氧钒(VOSO₄·5H₂O)、浓硫酸(98%)、乙酸(99%)，国药集团化学试剂有限公司。

1.2 UIO-66-NH₂的制备

UIO-66-NH₂合成原料摩尔比：ZrCl₄∶BDC-NH₂ = 1∶1。首先，将二者分别溶于20 mL DMF，再在配体溶液中加入2.3 mL乙酸，超声波辅助下混合20 min后得均匀的合成液。随后，将适量上述合成液倒入微波合成罐中，在不同温度下合成15 min，离心后获得棕黄色粉末。将上述粉末分散在无水乙醇中，多次更新乙醇下保持48 h，完成产物清洗。最后，在60℃下干燥12 h，获得产物UIO-66-NH₂，记为M-U66-NH₂。作为对比，采用同样的合成液，利用传统烘箱加热，在120℃条件下反应24 h制备UIO-66-NH₂，经过相同的清洗、干燥处理，获得产物UIO-66-NH₂，记为T-U66-NH₂。具体操作流程和样品编号见图1和表1。

1.3 复合质子交换膜的制备

以溶液浇筑法制备复合膜^[18]：称取一定质量的M-U66-NH₂，在超声辅助下分散在Nafion树脂溶液(固含量5%)中，在450 r/min下搅拌10 h后获得铸膜液。将适量铸膜液倒入带有凹槽的玻璃板上，在140℃真空烘箱中处理4 h，除去溶剂后得到复合膜，在恒温、恒湿条件下保存、备用。将掺杂M-U66-NH₂的复合膜记作M/N-X，X=1、2、3、6和9，X表示Nafion树脂中M-U66-NH₂的质量分数。将掺杂T-U66-NH₂的复合膜记作T/N-X，X=3。同时，采用相似工艺制备无M-U66-NH₂掺杂的纯树脂膜，记作P-N。具体操作流程和样品编号见图1和表1。

图 1 微波合成氨基官能化的UIO-66 (UIO-66-NH₂)及Nafion复合膜的制备过程示意图

Figure 1. Schematic diagram of the preparation process of UIO-66-NH₂ and Nafion composite membrane

M-U66-NH₂—UIO-66-NH₂ prepared by microwave assisted method; UIO-66-NH₂—Zr-MOF when the preparation method does not need to be distinguished; BDC-NH₂—2-aminoterephthalic Acid; DMF—N, N-dimethylformamide; HAc—Acetic acid

下载: 全尺寸图片幻灯片

表 1 复合质子交换膜的命名

Table 1. Naming of composite proton exchange membranes

Sample name	Instruction
M-U66-NH₂	"M" represents microwave heating; "U66" refers to the metal-organic framework material UIO-66; Overall, it indicates the UIO-66-NH₂ sample prepared by microwave heating.
T-U66-NH₂	"T" represents traditional oven heating; Overall, it indicates the UIO-66-NH₂ sample prepared by oven heating.
M/N-X	"M/N" represents the composite membrane with M-U66-NH₂ and Nafion resin; "X" represents the percentage content of M-U66-NH₂ in the membrane.
T/N-X	"T/N" represents the composite membrane with T-U66-NH₂ and Nafion resin; "X" represents the percentage content of T-U66-NH₂ added to the membrane.
P-N	A pure, unmodified Nafion membrane
N212	Commercial Nafion 212 membrane from DuPont (USA)

下载: 导出CSV

| 显示表格

1.4 形貌、结构表征和性能测试

采用X射线衍射仪(XRD，Bruker D8)和傅里叶红外光谱(FTIR，Nicolet 380，扫描范围：4000~750 cm⁻¹)评价粉末结构；扫描电子显微镜(SEM，SU8010，Hitachi)和扫描电子能谱(EDS)表征样品的微观形貌和元素分布；万能拉伸试验机(Instron 1186，加载速度为10 mm/min)检测膜的力学性能。

膜面积溶胀率的计算公式如下：

$\text{SR = }\frac{{({S}_{{\mathrm{wet}}}-{S}_{{\mathrm{dry}}}\text{)}}}{{{S}_{{\mathrm{dry}}}}}{\times 100\%}$

(1)

式中： ${S}_{{\mathrm{wet}}}$ 为湿膜面积(cm²)； ${S}_{{\mathrm{dry}}}$ 为干膜面积(cm²)。

吸水率的计算公式如下：

$\text{WU = }\frac{{({W}_{{\mathrm{wet}}}-{W}_{{\mathrm{dry}}}\text{)}}}{{{W}_{{\mathrm{dry}}}}}{\times 100\%}$

(2)

式中： ${W}_{{\mathrm{wet}}}$ 为湿膜质量(g)； ${W}_{{\mathrm{dry}}}$ 为干膜质量(g)。

应用紫外分光光度计(TU-1810，Beijing Purcell General Instrument Co., Ltd.)测试膜的VO²⁺渗透性^[7]，取样间隔12 h，钒离子渗透率计算公式如下：

${V}_{{\mathrm{B}}}\frac{{\mathrm{d}}\left({C}_{{\mathrm{B}}}\right(t\left)\right)}{{\mathrm{d}}t}=A\frac{P}{L}({C}_{{\mathrm{A}}}-{C}_{{\mathrm{B}}}(t\left)\right)$

(3)

式中： ${C}_{{\mathrm{A}}}$ 为VOSO₄溶液侧VO²⁺浓度； ${C}_{{\mathrm{B}}}\left(t\right)$ 为t时刻MgSO₄溶液侧VO²⁺浓度；V_B为MgSO₄溶液体积；P为钒离子渗透率；L为膜厚度；A为膜的有效面积(1.77 cm²)。

应用电化学工作站(CHI660E，上海辰华)测试膜的离子电导率( $\sigma$ )，计算公式如下：

$\sigma =\frac{{D}}{{{R}_{\text{b}}}}$

(4)

式中： ${R}_{{\mathrm{b}}}$ 为膜的面电阻；D为膜厚度。

采用新威电池测试系统表征VRB的电池性能，电流密度范围为100~200 mA·cm⁻²。循环测试的电流密度为150 mA·cm⁻²，循环200次。测试膜的放电容量衰减率，计算公式如下：

$容量衰减率=\frac{{Q_{{\mathrm{dis200}}}}}{{Q_{{\mathrm{dis1}}}}} {\times 100\%}$

(5)

式中：Q_dis200为循环200次后的放电容量；Q_dis1为第一次循环前放电容量。

$单次衰减=\frac{{容量衰减率}}{{循环次数}}$

(6)

2. 结果分析与讨论

2.1 UIO-66-NH₂的形貌与结构

为了优化UIO-66-NH₂的合成参数，考察微波合成温度对材料结构的影响。如图2(a)所示，合成时间为15 min，当温度较低时(60℃)，晶体结构较规则，颗粒尺寸约为1 μm，但产率仅为60%。随着合成温度的升高，晶体结构特征逐渐增强。如图2(c)所示，100℃下所合成样品具有明显的近八面体的晶体形貌^[18]。当温度升至120℃时，晶体结构更加完整，此时粒径约为200 nm (图2(d))，分散性较好，且产率接近90%。由图2(e)可见，通过普通水热法也可以制备出形貌规则的UIO-66-NH₂材料，只是所需的合成时间更长^[19]。

图 2 不同温度下微波加热15 min ((a)~(d))、烘箱加热24 h (e)制备UIO-66-NH₂的SEM图像及微波加热15 min样品的XRD图谱(f)

Figure 2. SEM images of UIO-66-NH₂ by microwave heating for 15 min ((a)-(d)), oven heating for 24 h (e) and XRD patterns of UIO-66-NH₂ by microwave heating for 15 min (f)

下载: 全尺寸图片幻灯片

XRD图谱反映样品的结晶度，如图2(f)所示，样品在2θ=7.40°和8.66°处出现了较明显的衍射峰，对应UIO-66-NH₂的(111)和(002)晶面，在2θ=26.22°和31.86°等处出现强度略低的衍射峰，以上峰位置和强度与文献报道结果基本一致^[20]，证明本实验成功合成出UIO-66-NH₂晶体。而且，随着合成温度的升高，晶体的衍射峰强度逐渐增加，即使在60℃下仍能合成出晶体。主要原因在于微波作用下反应釜内晶体前驱体溶液能够快速、均匀受热，微波的折射和反射诱导晶体快速成核、生长^[21]，合成效率得到显著提高。

2.2 质子交换膜的形貌与结构

为了更清晰地观察复合膜的微观形貌以及膜厚度，本实验对膜样品进行SEM表征。如图3(a)~3(c)所示，不同膜样品均呈现出较均匀、致密的形貌特征，M-U66-NH₂的掺杂量对膜形貌的影响较小。图3(d)~3(f)中，不同膜的厚度相近，约为40 μm。P-N膜的截面更加平滑，随着M-U66-NH₂掺杂量的增加，复合膜截面的粗糙度逐渐增加，但晶体的掺入未明显改变复合膜的结构。图3(g)~3(j)清晰地展现出M-U66-NH₂在M/N-3膜表面的元素分布情况，与M-U66-NH₂相关的Zr、N元素及与Nafion树脂相关的F、S元素在膜表面均匀分布，说明复合膜中M-U66-NH₂的分散性良好，未发生明显的团聚现象，这为功能材料充分发挥其表面性质和孔道尺寸等优势提供了条件。

图 3 P-N (a)、M/N-1 (b)、M/N-3 (c) 膜的表面SEM图像；P-N (d)、M/N-1 (e)、M/N-3 (f)膜的截面SEM图像；M/N-3膜的表面EDS元素分布图((g) N；(h) Zr；(i) F；(j) S)

Figure 3. Surface SEM images of P-N (a), M/N-1 (b) and M/N-3 (c); Cross-section SEM images of P-N (d), M/N-1 (e) and M/N-3 (f); EDS element images of M/N-3 ((g) N; (h) Zr; (i) F; (j) S)

下载: 全尺寸图片幻灯片

为了探究高M-U66-NH₂含量对复合膜结构的影响，本实验制备了M/N-6和M/N-9膜。如图4(a)所示，当掺杂量为6wt%时，复合膜截面出现较多尺寸约为2~3 μm的孔洞，这是由于较高掺杂量下，强相互作用促使M-U66-NH₂粒子团聚，导致膜内出现缺陷。当掺杂量增至9wt%时，复合膜截面呈现出非对称结构，膜层中可以观察到大尺寸的M-U66-NH₂团聚体。产生该现象的原因在于成膜过程中随着溶剂的缓慢挥发，较大尺寸的团聚体会逐渐下沉，并在玻璃板侧富集，形成高M-U66-NH₂含量的多孔、疏松区域，此结构不利于UIO-66-NH₂功能的充分发挥。同时，产生的缺陷会加速钒离子的跨膜渗透，降低质子交换膜的离子选择性和力学性能^[14]。

图 4 M/N-6 ((a), (c))和M/N-9 ((b), (d))膜的截面SEM图像

Figure 4. Cross-sectional SEM images of M/N-6 ((a), (c)) and M/N-9 ((b), (d))

下载: 全尺寸图片幻灯片

通过FTIR分析M-U66-NH₂、P-N、M/N-3膜样品的结构特征，如图5所示。M-U66-NH₂在3458、3356和1436 cm⁻¹处具有明显的特征峰，分别对应N—H的对称和非对称伸缩振动及N—H的剪切伸缩振动，与文献报道一致^[22]。P-N膜在1207、1151 cm⁻¹处为C—F伸缩振动峰，1053 cm⁻¹处为O=S=O振动峰及980 cm⁻¹处为C—F特征峰，此结果与文献中Nafion树脂的特征峰一致^[18]。M/N-3膜的FTIR图谱可以近似为M-U66-NH₂和P-N图谱的叠加，表明UIO-66-NH₂材料在膜中稳定存在，制膜过程并未破坏其微观结构。

图 5 M-U66-NH₂、P-N膜和M/N-3膜的红外图谱

Figure 5. FTIR spectra of M-U66-NH₂, P-N and M/N-3 membranes

下载: 全尺寸图片幻灯片

2.3 复合质子交换膜的理化性质

吸水率和溶胀率是质子交换膜的重要性能指标。如图6(a)所示，基于M-U66-NH₂的亲水性和丰富孔道，复合膜的吸水率和溶胀率均随掺入量的提高而增大^[23]。同时，M-U66-NH₂的刚性结构赋予膜较低的溶胀率(<4%)^[24]。膜的吸水率和溶胀率直接影响其离子传导性能，由图6(b)可见，随着UIO-66-NH₂掺杂量的提高，膜的质子传导率逐渐增大，而面电阻逐渐变小，如M/N-3的质子传导率达到122.18 mS·cm⁻¹，此改善效果得益于膜中—NH₂和—SO₃H形成的酸/碱对及存在的氢键网络^[25]。当掺杂量超过6wt%时，M-U66-NH₂在膜内团聚而缺陷产生，膜的质子传导率较低。比较可见，M/N-3膜的吸水率、溶胀率和质子传导率均高于T/N-3膜，这与UIO-66-NH₂的尺寸均一性和完整结构相关。

图 6 复合膜的吸水率和溶胀率(a)、面电阻和质子传导率(b)、力学性能(c)和应力-应变曲线(d)

Figure 6. Water uptake and swelling ratio (a), area resistance and conductivity (b), mechanical properties (c) and stress-strain curves of different membranes (d)

N212—Nafion 212 commercial membrane, DuPont, USA

下载: 全尺寸图片幻灯片

通过拉伸测试评价膜的机械强度。如图6(c)和图6(d)所示，M-U66-NH₂掺杂量低于3wt%时，复合膜的拉伸强度均高于P-N膜，如M/N-3膜的强度达到27 MPa，也高于T/N-3膜(22.3 MPa)。主要原因是UIO-66-NH₂具有刚性结构，且—NH₂基团与Nafion的—SO₃H基团可形成共轭酸碱对，其强相互作用促使Nafion的分子链与UIO-66-NH₂发生物理交联，从而提高了复合膜的力学性能^[26]。同时，尺寸更小、更规则的M-U66-NH₂与Nafion树脂间的分散更充分、交联程度更高，表现出更优的强化作用^[1]。当UIO-66-NH₂掺杂量超过6wt%时，膜中的孔洞缺陷导致其机械强度显著降低。

2.4 复合膜的钒离子渗透性和离子选择性

本实验选取的UIO-66-NH₂的有效孔径为0.52 nm，介于水分子(<0.3 nm)和钒离子(>0.6 nm)之间^[14]，可通过筛分效应提高膜的离子选择性。如图7(a)所示，M/N-X系列复合膜的阻钒性更好，如M/N-3膜的钒离子渗透浓度最低，仅为P-N膜的23%左右，而M/N-9膜的阻钒性能最差。当测试时间大于48 h，T/N-3膜的表现略差于M/N-3膜，主要原因是传统水热法所制备UIO-66-NH₂的晶体结构不够完善。

图 7 复合膜的钒离子渗透浓度(a)、钒离子渗透率和离子选择性(b)

Figure 7. Vanadium ion permeation concentration (a), vanadium ion permeability and ion selectivity (b) of composite membranes

下载: 全尺寸图片幻灯片

高钒离子渗透性不利于钒电池长期稳定运行，如图7(b)所示，优化条件下所制备的M/N-3膜的钒离子渗透率最低，仅为8.3×10⁻⁸ cm²·min⁻¹，且该膜的离子选择性达到15.6×10⁵ S·min·cm⁻³，约为P-N膜的30倍。当UIO-66-NH₂掺入量过低或过高时，膜的离子选择性均较低。另外，与M/N-3膜相比，掺杂T-U66-NH₂的复合膜具有接近的离子选择性，这说明两种方法所制备UIO-66-NH₂对复合膜离子选择性的影响较小。有研究结果指出增加厚度会提高膜的钒离子渗透率，进而降低膜的离子选择性^{[7, 27-29]}。因此，本实验在充分参考文献数据的基础上，为了平衡氢离子、钒离子的渗透情况，确定膜厚度为40 μm左右。

2.5 复合膜的电池性能

分别将P-N、M/N-X和T/N-3膜组装成模拟电池进行测试，评价不同电池的电压效率(Voltage efficiency，VE)、库伦效率(Coulombic efficiency，CE)和能量效率(EE)。由图8(a)可知，在相同电流密度下，随着M-U66-NH₂掺杂量的增加，电池的CE逐渐增大，在100 mA·cm⁻²电流密度下，M/N-3膜电池的库伦效率达到97.66%，高于P-N膜(95.8%)。因膜内粒子团聚造成了结构缺陷，M/N-6和M/N-9膜电池的CE低于M/N-3膜。图8(b)显示M/N-X复合膜的VE高于P-N膜，且随着电流密度的增加，复合膜所装配电池的VE逐渐减小，该趋势与文献报道相似^[5]。能量效率是电池性能的综合体现，由图8(c)可见，在电流密度为100~200 mA·cm⁻²的范围内，M/N-3膜的电池能量效率均高于P-N膜和N212膜(美国杜邦公司Nafion212商品膜)，最高达到83.8%，说明M-U66-NH₂的掺入有效提升了Nafion膜的电池性能。图8(d)~8(f)更清晰地显示了P-N膜、M/N-3膜和T/N-3膜的性能。在测试电流密度范围内，M/N-3膜的库伦效率与T/N-3膜相当，而其电压效率和能量效率均高于其他两种类型膜样品。可见，适宜比例M-U66-NH₂的引入有利于改善膜性能，进而提升钒电池的充放电效率。

图 8 不同膜所装配钒液流电池(VRBs)的库伦效率(CE) ((a), (d))、电压效率(VE) ((b), (e))和能量效率(EE) ((c), (f))

Figure 8. Coulombic efficiency (CE) ((a), (d)), voltage efficiency (VE) ((b), (e)) and energy efficiency (EE) ((c), (f)) of vanadium liquid flow battery (VRBs) assembled with different membranes

下载: 全尺寸图片幻灯片

在150 mA·cm⁻²电流密度下，对不同膜所装配电池充放电200次，比较其循环稳定性。如图9(a)所示，在200次的充放电测试过程中，M/N-3和T/N-3膜的电池CE、VE和EE相对较稳定，无明显降低，说明UIO-66-NH₂的引入有效阻碍了钒离子在膜内的渗透，保证了质子在膜中快速、稳定传输。此外，在测试周期内，M/N-3和T/N-3膜的能量效率分别稳定在77.8%和76.5%，无明显衰减，说明复合膜的理化稳定性可以满足钒电池工作环境的需求。但是，测试周期内不同电池的容量衰减情况差异较大。由图9(b)可见，P-N膜的电池容量衰减了81%，而经过同样测试后，掺杂UIO-66-NH₂膜的电池容量衰减率仅为45%左右，其中基于M-U66-NH₂的复合膜显示出更低的容量衰减率，单次衰减率仅为0.19%，较T/N-3膜(0.24%)提高0.05%，较P-N膜(0.41%)提高0.22%，同时，图9(c)证明了掺杂M-U66-NH₂的确会优化VRBs所用质子交换膜的电池性能。

图 9 150 mA·cm⁻²电流密度下复合膜所装配电池的循环效率(a)、容量保持率(b)以及与报道性能的对比(c)

Figure 9. Cycle efficiency (a) and capacity retention (b) of composite membranes at 150 mA·cm⁻², and comparisons with reported performance (c)

PBI—Polybenzimidazole; PS—Polystyrene; GO—Graphene oxide; SPEEK—Sulfonated poly(ether ether ketone); 2D-ZMs—Two-dimensional zeolite

下载: 全尺寸图片幻灯片

3. 结论

(1)与传统水热法相比，微波加热合成UIO-66-NH₂的效率更高，耗时仅为前者的1/96，且所制备晶体的结构更完整、更均匀，粒径约为200 nm，在Nafion溶液中分散性良好。

(2) UIO-66-NH₂可改善复合膜的理化稳定性，且可提高膜的质子传导性和离子选择性，优化条件下复合膜的质子传导率可达122.18 mS·cm⁻¹，离子选择性可达15.6×10⁵ S·min·cm⁻³。

(3)优化条件下复合膜体现出良好的电池性能。在150 mA·cm⁻²电流密度下，电池的能量效率大于77%，200次循环周期内单次容量衰减率为0.19%，较纯树脂膜提高0.22%。

(4)基于微波法的UIO-66-NH₂与Nafion形成的复合膜具有良好的理化性质和电池性能，为全钒液流电池用高性能质子交换膜的设计和制备提供了新的策略，具有良好的发展前景。

图 1 热防护系统与热防护材料^[1]

Figure 1. Thermal protection systems and materials^[1]

下载: 全尺寸图片幻灯片

图 2 ZrB₂-SiC高温氧化行为：(a)采用DNE-TGA分析ZrB₂-SiC在1600℃空气氧化不同时间后的材料微结构演化^[8]；(b)基于SEM与XEDS的ZrB₂-SiC在1627℃空气中氧化后扫描分析结果^[11]；(c)基于实验结果建立氧化产物演化模型示意图^[16]

Figure 2. High temperature oxidation behavior of ZrB₂-SiC: (a) DNE-TGA was used to analyze the microstructure evolution of ZrB₂-SiC after oxidation in air at 1600℃ for different times^[8]; (b) Scanning analysis results of ZrB₂-SiC after oxidation in air at 1627℃ based on SEM and XEDS^[11];(c) Schematic diagram of oxidation product evolution model based on experimental results^[16]

∆x—Thickness of oxide layer; t—Time; f_s—Volume fraction of SiC; I₁₂—Region 1-2; I₂₃—Region 2-3; I_3a—Region 3-a

下载: 全尺寸图片幻灯片

图 3 数据驱动的方法在陶瓷氧化损伤研究中的应用^[29-30]：(a)使用从文献中提取的实验数据库训练机器学习(ML)以预测氧化损伤； (b)模型预测氧化层厚度随温度变化和实际值的比较

UHTC—Ultra high temperature ceramics; T—Temperature; c—Composition; RD—Relative densification; GS—Grain size; d—Thickness of oxide layer

Figure 3. Application of data-driven methods in the study of ceramic oxidation damage^[29-30]: (a) Machine learning (ML) trained by the experimental database extracted from the literature to predict oxidative damage; (b) Comparison of model prediction and actual oxide scale thickness values vs. temperature

下载: 全尺寸图片幻灯片

图 4 碳纤维与C/SiC复合材料的高温氧化行为：(a) T300碳纤维的恒温氧化失重曲线^[31]；(b)平纹C/SiC复合材料的恒温氧化失重曲线^[31]；(c)基体裂纹提供氧气通道诱导局部纤维的氧化^[31]；(d)在裂纹尖端形成的碳纤维消耗区^[33]

Figure 4. High temperature oxidation behavior of carbon fiber and C/SiC composite: (a) Isothermal oxidation mass loss curves of T300 carbon fiber^[31];(b) Isothermal oxidation mass loss curves of plain C/SiC composite^[31]; (c) Matrix crack provides oxygen channel to induce local fiber oxidation^[31]; (d) Carbon fiber consumption zone formed at the crack tip^[33]

下载: 全尺寸图片幻灯片

图 5 C/ZrB₂-SiC复合材料的氧化失重^[36]：(a)试样的比例随温度的变化曲线；(b) 碳氧化失重和基体氧化增重的竞争

Δw∶S—Specific weight/Mass change per unit area

Figure 5. Oxidation mass loss of C/ZrB₂-SiC composites^[36]: (a) Curve of specific gravity of samples with temperature; (b) Competition between carbon oxidation mass loss and matrix oxidation mass gain

下载: 全尺寸图片幻灯片

图 6 C/ZrB₂-SiC复合材料的高温氧化行为：(a)不同深度的氧化层微结构成分与形貌^[36]；(b)轴向纤维和横向纤维在空气中的氧化行为示意图^[34]

Figure 6. High temperature oxidation behavior of C/ZrB₂-SiC composites: (a) Microstructure composition and morphology of oxide layers at different depths^[36]; (c) Oxidation behavior of axial fibers and transverse fibers in air^[34]

下载: 全尺寸图片幻灯片

图 7 C/SiC复合材料的裂纹闭合行为^[32]：((a), (b))高温氧化后基体裂纹弥合；(c)裂纹闭合导致性能提升

Figure 7. Crack closure behavior of C/SiC composites^[32]: ((a), (b)) Crack closure in SiC matrix after high temperature oxidation; (c) Crack closure leads to improved performance

下载: 全尺寸图片幻灯片

图 8 C/SiC复合材料的应力氧化行为^[45]：(a)均匀与非均匀氧化共存形貌(归一化应力NS=0.32)；(b)非均匀氧化形貌(归一化应力NS=0.64)；(c) 在不同应力氧化机制下C/SiC复合材料的长度变化

Figure 8. Stress oxidation behavior of C/SiC composites^[45]: (a) Uniform/non-uniform fiber oxidation coexistence (Normalized stress NS=0.32); (b) Non-uniform fiber oxidation (NS=0.64); (c) Length changes of C/SiC composites under different stress oxidation mechanisms

下载: 全尺寸图片幻灯片

图 9 C/SiC复合材料的耦合失效行为：(a)拉伸载荷作用下裂纹的演化与局部纤维结构相关^[47]；((b), (c))基于原位Micro-CT表征SiC/SiC的应力氧化力学行为和结构特征^[53]

Figure 9. Coupling failure behavior of ceramic matrix composites: (a) Evolution of cracks under tensile load is related to local fiber structure^[47]; ((b), (c)) Characterization of stress oxidation mechanical behavior and structural characteristics of SiC/SiC based on in-situ Micro-CT^[53]

下载: 全尺寸图片幻灯片

图 10 陶瓷基复合材料冲击损伤实验以及剩余性能研究：(a)氧乙炔对冲击后C/C复合材料和SiC-C/C复合材料氧化测试与损伤表征^[57]；(b)氧化质量损失与涂层冲击损伤面积的关系以及C/SiC复合材料剩余弯曲性能^[58]

v—Velocity

Figure 10. Impact damage experiment and residual performance study of ceramic matrix composites: (a) Oxidation test and damage characterization of oxyacetylene on C/C composites and SiC-C/C composites after impact^[57]; (b) Relationship between oxidation mass loss and impact damage area of coating and C/SiC residual bending properties of composites^[58]

下载: 全尺寸图片幻灯片

图 11 陶瓷基复合材料的宏细观氧化模型：(a)基于不同控制因素的宏观氧化模型^[60]；(b)建立考虑预制裂纹的纤维横向氧化细观有限元模型^[62]；(c)基于剪滞理论的细观氧化损伤模型^[65]

σ—Stress; V_f—Fiber volume fraction; τ_i—Interface shear force in slip region; τ_f—Interface shear force in the oxidized region

Figure 11. Macro and meso scale oxidation models of ceramic matrix composites: (a) Macro oxidation model based on different control factors^[60]; (b) Meso finite element model of fiber transverse oxidation considering prefabricated cracks^[62]; (c) Meso oxidation damage model based on shear lag theory^[65]

下载: 全尺寸图片幻灯片

图 12 陶瓷基复合材料的耦合失效模型：(a)非线性宏-微观耦合二维有限元模型^[68]；(b)氧气浓度计算结果^[68]；(c)基于Micro-CT建立的细观单胞模型^[26]；(d)基于用户子程序的氧化损伤模拟云图^[26]；(e)结合氧化动力学和渐近损伤模型的失效分析方法^[71]

Figure 12. Coupling failure model of ceramic matrix composites: (a) Nonlinear macro micro coupled two-dimensional finite element model^[68]; (b) Oxygen concentration calculation^[68]; (c) Meso model based on Micro-CT^[26]; (d) Oxidation damage cloud figure simulated by user subroutine^[26]; (e) Failure analysis method combining oxidation kinetics and asymptotic damage models^[71]

${{P}_{{{\text{O}}_{\text{2}}}}}$ —Oxygen partial pressure; $\nabla$ X—Displacement gradient; $\nabla$ ρ—Density gradient; $\nabla$ T—Temperature gradient; C_ijkl—Stiffness coefficient; K_ij—Thermal conductivity; D_ij—Diffusion coefficient; q—Net flux; n—External normal direction; S, S^I —Outer and inner boundaries; V⁻—Volume occupied by fibers; V⁺—Volume occupied by materials other than fibers; p—Length ratio; η—Number ratio

下载: 全尺寸图片幻灯片

参考文献(71)

[1]	UYANNA O, NAJAFI H. Thermal protection systems for space vehicles: A review on technology development, current challenges and future prospects[J]. Acta Astronautica, 2020, 176: 341-356. DOI: 10.1016/j.actaastro.2020.06.047
[2]	SAIROCZAK D, SMITH H. A review of design issues specific to hypersonic flight vehicles[J]. Progress in Aerospace Sciences, 2016, 84: 1-28. DOI: 10.1016/j.paerosci.2016.04.001
[3]	关春龙, 李垚, 赫晓东. 可重复使用热防护系统防热结构及材料的研究现状[J]. 宇航材料工艺, 2003, 33(6): 7-11. GUAN Chunlong, LI Yao, HE Xiaodong. Research status of structures and materials for reusable TPS[J]. Aerospace Materials & Technology, 2003, 33(6): 7-11(in Chinese).
[4]	包为民. 可重复使用运载火箭技术发展综述[J]. 航空学报, 2023, 44(23): 629555. BAO Weimin. A review of reusable launch vehicle technology development[J]. Acta Aeronautica et Astronautica Sinica, 2023, 44(23): 629555(in Chinese).
[5]	国义军, 石卫波, 曾磊, 等. 高超声速飞行器烧蚀防热理论与应用[M]. 北京: 科学出版社, 2019: 12. GUO Yijun, SHI Weibo, ZENG Lei, et al. Mechanism of ablative thermal protection applied to hypersonic vehicles[M]. Beijing: China Science Publishing & Media Ltd., 2019: 12(in Chinese).
[6]	SHVYDYUK K O, NUNES-PEREIRA J, RODRIGUES F F, et al. Review of ceramic composites in aeronautics and aerospace: A multifunctional approach for TPS, TBC and DBD applications[J]. Ceramics, 2023, 6(1): 195-230. DOI: 10.3390/ceramics6010012
[7]	WEI F, ZHANG Y, LEI L, et al. High-frequent pulsing ablation of C/C-SiC-ZrB₂-ZrC composite for different cycles to 2000 times in plasma[J]. Transactions of Nonferrous Metals Society of China, 2023, 33(2): 553-562. DOI: 10.1016/S1003-6326(22)66127-2
[8]	MILLER-OANA M, CORRAL E L. High-temperature isothermal oxidation of ultra-high temperature ceramics using thermal gravimetric analysis[J]. Journal of the American Ceramic Society, 2016, 99(2): 619-626. DOI: 10.1111/jace.14001
[9]	GRIGORIEV O N, STEPANENKO A V, VINOKUROV V B, et al. ZrB₂-SiC ceramics: Residual stresses and mechanical properties[J]. Journal of the European Ceramic Society, 2021, 41(9): 4720-4727. DOI: 10.1016/j.jeurceramsoc.2021.02.053
[10]	TIWARI M, SINGH V K. Influence of SiC content to control morphology of in-situ synthesized ZrB₂-SiC composite through single-step reduction process[J]. Vacuum, 2022, 203: 111251. DOI: 10.1016/j.vacuum.2022.111251
[11]	LEVINE S R, OPILA E J, HALBIG M C, et al. Evaluation of ultra-high temperature ceramics for aeropropulsion use[J]. Journal of the European Ceramic Society, 2002, 22(14-15): 2757-2767. DOI: 10.1016/S0955-2219(02)00140-1
[12]	CHO Y J, LU K. High temperature oxidation behaviors of bulk SiC with low partial pressures of air and water vapor in argon[J]. Corrosion Science, 2020, 174: 108795. DOI: 10.1016/j.corsci.2020.108795
[13]	DEAL B E, GROVE A S. General relationship for the thermal oxidation of silicon[J]. Journal of Applied Physics, 1965, 36(12): 3770-3778. DOI: 10.1063/1.1713945
[14]	RAJ R. Chemical potential-based analysis for the oxidation kinetics of Si and SiC single crystals[J]. Journal of the American Ceramic Society, 2013, 96(9): 2926-2934. DOI: 10.1111/jace.12464
[15]	FAHRENHOLTZ W G. Thermodynamic analysis of ZrB₂-SiC oxidation: Formation of a SiC-depleted region[J]. Journal of the American Ceramic Society, 2007, 90(1): 143-148. DOI: 10.1111/j.1551-2916.2006.01329.x
[16]	PARTHASARATHY T A, RAPP R A, OPEKA M, et al. Modeling oxidation kinetics of SiC-containing refractory diborides[J]. Journal of the American Ceramic Society, 2012, 95(1): 338-349. DOI: 10.1111/j.1551-2916.2011.04927.x
[17]	PARTHASARATHY T A, RAPP R A, OPEKA M, et al. A model for the oxidation of ZrB₂, HfB₂ and TiB₂[J]. Acta Materialia, 2007, 55(17): 5999-6010. DOI: 10.1016/j.actamat.2007.07.027
[18]	MA Y, YAO X, HAO W, et al. Oxidation mechanism of ZrB₂/SiC ceramics based on phase-field model[J]. Composites Science and Technology, 2012, 72(10): 1196-1202. DOI: 10.1016/j.compscitech.2012.04.003
[19]	CHEN X, SUN Z, CHEN Z, et al. ReaxFF molecular dynamics simulation of oxidation behavior of 3C-SiC in O₂ and CO₂[J]. Computational Materials Science, 2021, 191: 110341. DOI: 10.1016/j.commatsci.2021.110341
[20]	ZHANG P, ZHANG Y, CHEN G, et al. High-temperature oxidation behavior of CVD-SiC ceramic coating in wet oxygen and structural evolution of oxidation product: Experiment and first-principle calculations[J]. Applied Surface Science, 2021, 556: 149808. DOI: 10.1016/j.apsusc.2021.149808
[21]	ZUMPICCHIAT G, PASCAL S, TUPIN M, et al. Finite element modelling of the oxidation kinetics of Zircaloy-4 with a controlled metal-oxide interface and the influence of growth stress[J]. Corrosion Science, 2015, 100: 209-221. DOI: 10.1016/j.corsci.2015.07.024
[22]	DONG X, FANG X, FENG X, et al. Diffusion and stress coupling effect during oxidation at high temperature[J]. Journal of the American Ceramic Society, 2013, 96(1): 44-46. DOI: 10.1111/jace.12105
[23]	LOEFFEL K, ANAND L. A chemo-thermo-mechanically coupled theory for elastic-viscoplastic deformation, diffusion, and volumetric swelling due to a chemical reaction[J]. International Journal of Plasticity, 2011, 27(9): 1409-1431. DOI: 10.1016/j.ijplas.2011.04.001
[24]	CHEN L, YUEMING L. A coupled mechanical-chemical model for reflecting the influence of stress on oxidation reactions in thermal barrier coating[J]. Journal of Applied Physics, 2018, 123(21): 215305. DOI: 10.1063/1.5019848
[25]	ZHOU Z, PENG X, WEI Z. A thermo-chemo-mechanical model for the oxidation of zirconium diboride[J]. Journal of the American Ceramic Society, 2015, 98(2): 629-636. DOI: 10.1111/jace.13333
[26]	ZHAO Y, CHEN Y, HE C, et al. A damage-induced short-circuit diffusion model applied to the oxidation calculation of ceramic matrix composites (CMCs)[J]. Composites Part A: Applied Science and Manufacturing, 2019, 127: 105621. DOI: 10.1016/j.compositesa.2019.105621
[27]	WANG H, SHEN S. A chemomechanical coupling model for oxidation and stress evolution in ZrB₂-SiC[J]. Journal of Materials Research, 2017, 32(7): 1267-1278. DOI: 10.1557/jmr.2016.489
[28]	RULKO T A, PICKARD D N, RADOVITZKY R. Fully-coupled multiphysics simulations of the thermo-mechanical oxidation of ceramics for spacecraft heat shields[C]//AIAA SCITECH 2024 Forum. Reston, VA: American Institute of Aeronautics and Astronautics, 2024: 0029.
[29]	BIANCO G, NISAR A, ZHANG C, et al. Predicting oxidation damage in ultra high-temperature borides: A machine learning approach[J]. Ceramics International, 2022, 48(20): 29763-29769. DOI: 10.1016/j.ceramint.2022.06.236
[30]	BIANCO G, NISAR A, ZHANG C, et al. Predicting oxidation damage of ultra high-temperature carbide ceramics in extreme environments using machine learning[J]. Ceramics International, 2023, 49(12): 19974-19981. DOI: 10.1016/j.ceramint.2023.03.119
[31]	HALBIG M C, MCGUFFIN-CAWLEY J D, ECKEL A J, et al. Oxidation kinetics and stress effects for the oxidation of continuous carbon fibers within a microcracked C/SiC ceramic matrix composite[J]. Journal of the American ceramic society, 2008, 91(2): 519-526. DOI: 10.1111/j.1551-2916.2007.02170.x
[32]	OPILA E J, SERRA J L. Oxidation of carbon fiber-reinforced silicon carbide matrix composites at reduced oxygen partial pressures[J]. Journal of the American Ceramic Society, 2011, 94(7): 2185-2192. DOI: 10.1111/j.1551-2916.2010.04376.x
[33]	JACOBSON N S, CURRY D M. Oxidation microstructure studies of reinforced carbon/carbon[J]. Carbon, 2006, 44(7): 1142-1150. DOI: 10.1016/j.carbon.2005.11.013
[34]	VINCI A, ZOLI L, SCITI D. Influence of SiC content on the oxidation of carbon fibre reinforced ZrB₂/SiC composites at 1500 and 1650°C in air[J]. Journal of the European Ceramic Society, 2018, 38(11): 3767-3776. DOI: 10.1016/j.jeurceramsoc.2018.04.064
[35]	VINCI A, ZOLI L, LANDI E, et al. Oxidation behaviour of a continuous carbon fibre reinforced ZrB₂-SiC composite[J]. Corrosion Science, 2017, 123: 129-138. DOI: 10.1016/j.corsci.2017.04.012
[36]	ZHANG D, HU P, DONG S, et al. Oxidation behavior and ablation mechanism of C_f/ZrB₂-SiC composite fabricated by vibration-assisted slurry impregnation combined with low-temperature hot pressing[J]. Corrosion Science, 2019, 161: 108181. DOI: 10.1016/j.corsci.2019.108181
[37]	ZOLI L, SCITI D. Efficacy of a ZrB₂-SiC matrix in protecting C fibres from oxidation in novel UHTCMC materials[J]. Materials & Design, 2017, 113: 207-213.
[38]	VINCI A, REIMER T, ZOLI L, et al. Influence of pressure on the oxidation resistance of carbon fiber reinforced ZrB₂/SiC composites at 2000 and 2200℃[J]. Corrosion Science, 2021, 184: 109377. DOI: 10.1016/j.corsci.2021.109377
[39]	ZHAO Z, LI K, LI W. Ablation behavior of ZrC-SiC-ZrB₂ and ZrC-SiC inhibited carbon/carbon composites components under ultrahigh temperature conditions[J]. Corrosion Science, 2021, 189: 109598. DOI: 10.1016/j.corsci.2021.109598
[40]	KOU S, MA J, MA Y, et al. Microstructure and flexural strength of C/HfC-ZrC-SiC composites prepared by reactive melt infiltration method[J]. Journal of the European Ceramic Society, 2023, 43(5): 1864-1873. DOI: 10.1016/j.jeurceramsoc.2022.12.015
[41]	ZHANG Y, ZHANG L, CHENG L, et al. Tensile behavior and microstructural evolution of a carbon/silicon carbide composite in simulated re-entry environments[J]. Materials Science and Engineering: A, 2008, 473(1-2): 111-118. DOI: 10.1016/j.msea.2007.05.015
[42]	ZHANG Y, ZHANG L, LIU Y, et al. Oxidation effects on in-plane and interlaminar shear strengths of two-dimensional carbon fiber reinforced silicon carbide composites[J]. Carbon, 2016, 98: 144-156. DOI: 10.1016/j.carbon.2015.10.091
[43]	CHENG T, WANG X, ZHANG R, et al. Tensile properties of two-dimensional carbon fiber reinforced silicon carbide composites at temperatures up to 2300°C[J]. Journal of the European Ceramic Society, 2020, 40(3): 630-635. DOI: 10.1016/j.jeurceramsoc.2019.10.030
[44]	CHENG T. Insights into fracture mechanisms and strength behaviors of two-dimensional carbon fiber reinforced silicon carbide composites at elevated temperatures[J]. Journal of the European Ceramic Society, 2022, 42(1): 71-86. DOI: 10.1016/j.jeurceramsoc.2021.09.028
[45]	LUAN X, CHENG L, ZHANG J, et al. Effects of temperature and stress on the oxidation behavior of a 3D C/SiC composite in a combustion wind tunnel[J]. Composites Science and Technology, 2010, 70(4): 678-684. DOI: 10.1016/j.compscitech.2009.12.025
[46]	MORSCHER G N. Stress-dependent matrix cracking in 2D woven SiC-fiber reinforced melt-infiltrated SiC matrix composites[J]. Composites Science and Technology, 2004, 64(9): 1311-1319. DOI: 10.1016/j.compscitech.2003.10.022
[47]	SEVEVNER K M, TRACY J M, ZHE C, et al. Crack opening behavior in ceramic matrix composites[J]. Journal of the American Ceramic Society, 2017, 100(10): 4734-4747.
[48]	FENG Y, FENG Z, LI S, et al. Micro-CT characterization on porosity structure of 3D C_f/SiC_m composite[J]. Composites Part A: Applied Science and Manufacturing, 2011, 42(11): 1645-1650. DOI: 10.1016/j.compositesa.2011.07.015
[49]	CHATEAU C, GELEBART L, BORNERT M, et al. In situ X-ray microtomography characterization of damage in SiC_f/SiC minicomposites[J]. Composites Science and Technology, 2011, 71(6): 916-924. DOI: 10.1016/j.compscitech.2011.02.008
[50]	GOULMY J P, CATY O, REBILLAT F. Characterization of the oxidation of C/C/SiC composites by X-ray micro-tomography[J]. Journal of the European Ceramic Society, 2020, 40(15): 5120-5131.
[51]	AI S, SONG W, CHEN Y. Stress field and damage evolution in C/SiC woven composites: Image-based finite element analysis and in situ X-ray computed tomography tests[J]. Journal of the European Ceramic Society, 2021, 41(4): 2323-2334. DOI: 10.1016/j.jeurceramsoc.2020.12.026
[52]	MAZARS V, CATY O, COUEGNAT G, et al. Damage investigation and modeling of 3D woven ceramic matrix composites from X-ray tomography in-situ tensile tests[J]. Acta Materialia, 2017, 140: 130-139. DOI: 10.1016/j.actamat.2017.08.034
[53]	BALE H A, HABOUB A, MACDOWELL A A, et al. Real-time quantitative imaging of failure events in materials under load at temperatures above 1600℃[J]. Nature Materials, 2013, 12(1): 40-46. DOI: 10.1038/nmat3497
[54]	CHENG T, ZHANG R, PEI Y, et al. Tensile properties of two-dimensional carbon fiber reinforced silicon carbide composites at temperatures up to 1800℃ in air[J]. Extreme Mechanics Letters, 2019, 31: 100546. DOI: 10.1016/j.eml.2019.100546
[55]	LEISER D B, CHURCHWARD R, KATVALA V, et al. Advanced porous coating for low-density ceramic insulation materials[J]. Journal of the American Ceramic Society, 1989, 72(6): 1003-1010. DOI: 10.1111/j.1151-2916.1989.tb06259.x
[56]	GROSCH D, BERTRAND F. Thermal protection system (TPS) impact experiments[C]//47th AIAA/ASME/ASCE/AHS/ASC Structures, Structural Dynamics, and Materials Conference. Newport, Rhode Island: AIAA, 2006: 1780.
[57]	XUE L, LI K, JIA Y, et al. Effects of hypervelocity impact on ablation behavior of SiC coated C/C composites[J]. Materials & Design, 2016, 108: 151-156.
[58]	YAO L, LYU P, BAI G, et al. Influence of low velocity impact on oxidation performance of SiC coated C/SiC composites[J]. Ceramics International, 2019, 45(16): 20470-20477. DOI: 10.1016/j.ceramint.2019.07.025
[59]	SULLIVAN R M. A model for the oxidation of carbon silicon carbide composite structures[J]. Carbon, 2005, 43(2): 275-285. DOI: 10.1016/j.carbon.2004.09.010
[60]	HALBIG M C. Stressed oxidation and modeling of C/SiC in oxidizing environments[C]//25th Annual Conference on Composites, Advanced Ceramics, Materials, and Structures: A: Ceramic Engineering and Science Proceedings. Hoboken: John Wiley & Sons, Inc., 2001: 625-632.
[61]	LAMOUROUX F, NASLAIN R, JOUIN J M. Kinetics and mechanisms of oxidation of 2D woven C/SiC composites: II, theoretical approach[J]. Journal of the American Ceramic Society, 1994, 77(8): 2058-2068. DOI: 10.1111/j.1151-2916.1994.tb07097.x
[62]	XU Y, ZHANG P, LU H, et al. Numerical modeling of oxidized C/SiC microcomposite in air oxidizing environments below 800℃: Microstructure and mechanical behavior[J]. Journal of the European Ceramic Society, 2015, 35(13): 3401-3409. DOI: 10.1016/j.jeurceramsoc.2015.05.039
[63]	SUN Z, NIU X, WANG Z, et al. Verification and prediction of residual strength of C/SiC composites under non-stress oxidation[J]. Journal of Materials Science, 2014, 49(23): 8192-8203. DOI: 10.1007/s10853-014-8528-1
[64]	CASAS L, MARTINEZ-ESNAOLA J M. Modelling the effect of oxidation on the creep behaviour of fibre-reinforced ceramic matrix composites[J]. Acta Materialia, 2003, 51(13): 3745-3757. DOI: 10.1016/S1359-6454(03)00189-7
[65]	LI L. Modeling matrix cracking of fiber-reinforced ceramic-matrix composites under oxidation environment at elevated temperature[J]. Theoretical and Applied Fracture Mechanics, 2017, 87: 110-119. DOI: 10.1016/j.tafmec.2016.11.003
[66]	DENG Y, LI W, MA J, et al. Thermal-mechanical-oxidation coupled first matrix cracking stress model for fiber reinforced ceramic-matrix composites[J]. Journal of the European Ceramic Society, 2021, 41(7): 4016-4024. DOI: 10.1016/j.jeurceramsoc.2021.02.033
[67]	LEE S, SUNDARATAGHAVAN V. Multi-scale modeling of moving interface problems with flux and field jumps: Application to oxidative degradation of ceramic matrix composites[J]. International Journal for Numerical Methods in Engineering, 2011, 85(6): 784-804. DOI: 10.1002/nme.2996
[68]	MEI H, CHENG L, ZHANG L, et al. Modeling the effects of thermal and mechanical load cycling on a C/SiC composite in oxygen/argon mixtures[J]. Carbon, 2007, 45(11): 2195-2204. DOI: 10.1016/j.carbon.2007.06.051
[69]	LUAN X, CHENG L, XIE C. Stressed oxidation life predication of 3D C/SiC composites in a combustion wind tunnel[J]. Composites Science and Technology, 2013, 88: 178-183. DOI: 10.1016/j.compscitech.2013.09.001
[70]	CHEN X, SUN Z, LI H, et al. Modeling the effect of oxidation on the residual tensile strength of SiC/C/SiC mini composites in stressed oxidizing environments[J]. Journal of Materials Science, 2020, 55(8): 3388-3407. DOI: 10.1007/s10853-019-04255-4
[71]	SANTHOSH U, AHMAD J, OIARD G, et al. A synergistic model of stress and oxidation induced damage and failure in silicon carbide-based ceramic matrix composites[J]. Journal of the American Ceramic Society, 2021, 104: 4163-4182.

施引文献

资源附件(0)

长摘要

目的
碳纤维增韧陶瓷基复合材料兼具陶瓷材料优良的抗氧化腐蚀性能和碳纤维材料增强增韧的力学性能，已成为最有潜力的高超声速飞行器热防护候选材料。碳纤维增韧陶瓷基复合材料在多物理场耦合服役环境下的高温氧化损伤失效机理对热防护材料设计及性能表征与评价至关重要，也一直是国内外学者研究的热点。本文从高温氧化机理、耦合失效实验、高温氧化模型三个方面对C/SiC和C/ZrB-SiC复合材料进行详细论述、总结和展望，以期为碳纤维增韧陶瓷基复合材料在热力氧耦合条件下的热/力响应及性能评价研究奠定理论模型基础。
方法
(1)在高温氧化机理研究方面，主要是采用非等温和等温静态氧化测试以及动态氧化测试分析SiC基陶瓷及其复合材料的氧化动力学行为，通过热重分析(TGA)、微结构观测(SEM)以及物相分析(XRD)研究不同氧化条件下材料的宏微观物理状态以及化学成分的演化。(2)在氧化损伤力学实验研究方面，主要是开展无应力/应力氧化的高温力学实验、氧化后低速/高速冲击实验，获得氧化后复合材料的剩余力学性能，揭示氧化损伤和外界环境因素关联性。(3)在氧化损伤模型研究方面，基于多孔介质理论和质量守恒方程建立了氧化模型，通过引入组分材料的氧化动力学建模分析复合材料的复杂氧化行为。借助有限元软件对具有复杂结构的复合材料的氧化损伤过程。
结果
(1)在高温氧化机理研究方面，获得了SiC组分主、被动氧化机制以及ZrB-SiC的全温区非烧蚀的形成机理。陶瓷基复合材料内部的孔隙、裂纹以及纤维氧化后的孔洞成为氧气向内部运输的通道。随着温度和氧分压的变化，存在反应和扩散控制的这两种不同的氧化主控因素。陶瓷基复合材料的氧化与温度和局部结构相关。(2)在氧化损伤力学实验研究方面，氧化损伤累积会弱化裂纹偏转、界面脱粘以及纤维拔出的增韧机制。氧化后碳纤维增韧陶瓷基复合材料的剩余强度具有温度依赖性，氧化在“中等高”温度条件下是最具破坏性的。局部的拉伸应力会延迟裂纹闭合的时间和促进纤维的暴露氧化，进而改变裂纹扩散控制的氧化机制。(3)在氧化损伤模型研究方面。宏观模型，主要是通过多孔介质理论和质量守恒方程建立了均质连续氧化模型，不能解释碳纤维增韧陶瓷基复合材料局部的氧化损伤行为。微细观模型，可以采用适当的均质化技术实现尺度之间的信息传递。引入氧化动力学建模后可以分析局部氧化行为和氧化后的剩余力学性能。多尺度模型，可以同时模拟出均匀化尺度氧气浓度分布和组分材料尺度氧化损伤程度，但缺乏考虑孔隙演化与纤维氧化的相互作用机制，也缺乏考虑基体氧化和载荷作用下基体损伤演化对气体扩散的影响。
结论
本文以C/SiC和C/ZrB-SiC复合材料两种典型的碳纤维增韧陶瓷基复合材料为对象，详细综述、归纳和分析了该类复合材料的高温氧化机理、氧化损伤力学实验和氧化损伤模型三方面的研究工作，并展望了碳纤维增韧陶瓷基复合材料氧化损伤研究的发展趋势。