Mo_f/TiAl复合材料的制备及变形断裂行为研究
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摘要
高温结构材料是航空航天推进系统实现革命性变革与发展的关键因素。TiAl基合金因具有密度低、强度高、抗氧化性能好及抗蠕变性能优异等优点成为非常有竞争力的高温结构材料，目前已应用于发动机叶片、涡轮与排气阀等构件。但航空航天工业的迅猛发展，对发动机材料提出了更高的要求。TiAl基合金要想获得更广泛的应用并满足更为苛刻的服役条件，需要进一步改善其室温断裂韧性，并提高高温强度。本文通过采用连续Mo纤维增强的方式来实现TiAl基合金强韧化的目的。采用了粉末浆料铸造与真空热压烧结相结合的工艺方法成功地制备了Mo_f/TiAl复合材料，研究了制备工艺参数对复合材料显微组织与力学性能影响，观察分析了复合材料的变形与断裂行为，揭示了复合材料的强韧化机理。
本文采用改进的粉末浆料铸造法简便有效地实现了纤维与基体粉末的预复合。选用PMMA、丙酮与Ti、Al混合粉末制备粉末浆料，并在三者配比为10g:70ml:60g时，成功地获得了致密无空洞、纤维排布均匀、Ti、Al粉末分散良好的品质优良的预制体。预制体经除气与热压烧结两步最终制得复合材料。
制备的复合材料中纤维排布均匀，保持了制备前的高长径比的纤维组织形貌；基体以TiAl相为主，并含有少量的Ti3Al相。热压烧结温度升高与时间延长，有利于基体致密化程度及组织均匀性的提高，但加剧界面反应，使界面处生成两个连续的反应层，由Mo纤维到基体依次为呈现放射型柱状晶形貌的δ-(Mo,Ti)3Al相和粗大等轴晶形貌的β'-(Mo,Ti)A1相。纳米压痕测试表明，δ相的硬度及弹性模量高于β'。纤维推出试验测得的界面剪切强度大于367MPa，Mo/δ界面发生脱粘。界面相的生长动力学研究表明，δ相与β'相的生长均遵循抛物线规律，二者的生长速度随温度的升高而加快。
复合材料的性能受制备工艺参数影响，最优工艺参数为380℃保温除气1h，并经1100℃热压烧结1h。制备的单向纤维增强的复合材料的室温纵向与横向弯曲强度分别为735.4MPa与249.3MPa，纵向弯曲强度受加载方向影响很小。正交纤维增强的复合材料的室温弯曲强度下降至374.4MPa。700℃-1000℃范围内，单向纤维与正交纤维增强的复合材料的弯曲强度均随温度的升高先增加后降低，二者均在800℃时达到最高值，分别为762.9MPa与564.4MPa。
复合材料的压缩屈服强度随测试温度的升高而减低；纵向压缩屈服强度比法向略高，纤维排布方式对复合材料的法向压缩强度影响较小。单边缺口梁法测得的单向纤维增强的复合材料的断裂韧性为23.55MPa·m1/2，较基体TiAl提高50%以上。
由于脆性基体与韧性连续纤维的共同作用，Mo_f/TiAl复合材料的变形与断裂过程不同于均质材料，变形与断裂过程相互交织，其间伴随着基体与纤维承载作用的变化。室温纵向弯曲与拉伸时，复合材料的变形断裂过程分为三个阶段：第一阶段，复合材料的整体变形来自于纤维与基体间协同发生的弹性变形，纤维与基体共同承载；第二阶段，基体与纤维变形失配，脆性基体通过裂纹的不断萌生扩展匹配复合材料的变形，其有效承载下降，纤维则继续发生弹性变形，承载不断增加；第三阶段，基体因裂纹数量饱和而完全失效，纤维独自承载，复合材料的变形完全由纤维的弹、塑性变形提供，并最终在纤维颈缩断裂后复合材料完全破坏。复合材料中裂纹萌生于基体，发生沿TiAl相晶界及穿过Ti3Al相晶粒的扩展后，直接穿过δ与β'两界面相而受阻于纤维，两界面相与基体一同变形开裂。Mo/δ、δ/β'与β'/TiAl三界面在纤维塑性变形前均完好无损，之后因脆性δ相与韧性纤维难以协调变形，Mo/δ界面脱粘。Mo纤维对裂纹不敏感，基体裂纹难以扩展穿过纤维，纤维在达到自身的抗拉强度后萌生裂纹并断裂。
高温下复合材料的弯曲变形与断裂方式与室温时不同，塑韧性大为提高的基体能够通过较大的变形而非裂纹的萌生扩展来匹配复合材料的整体变形，在纤维发生颈缩断裂后，复合材料因基体承载能力的不足而快速断裂。
纵向压缩载荷作用下，复合材料中纤维在与纵向成45°方向的最大切应力作用下发生较大弯曲而不断裂，其间基体则沿纤维轴线破碎，并造成界面Mo/δ开裂。法向压缩载荷作用下，复合材料的剪切破坏面平行于纤维轴向，并呈现台阶状。高温时，复合材料塑韧性增加，基体与纤维变形协调一致，复合材料发生墩粗变形而不破裂。
通过对复合材料力学性能及变形与断裂行为的综合分析，揭示了复合材料的强韧化机理：较高的基体致密度、良好的界面结合及较少的纤维损耗三者之间的平衡是复合材料强度得到最大提高的保证；纤维的桥接、塑性变形、脱粘与拔出以及裂纹的偏转枝生是复合材料的主要韧化机制，其中，纤维的塑性变形使强韧性优异的Mo纤维的增韧效果远胜于脆性陶瓷纤维。
High-temperature structural materials play a key role on the realization of therevolutionary development of the aerospace propulsion system. Among them,TiAl-based alloys have become very competitive ones, because they have theadvantages of low density, high strength, good oxidation resistance and excellentcreep resistance. Currently, they have been applied to engine blades, turbine,exhaust valves and other components. However, the rapid development of theaerospace industry proposes higher demands on engine materials. In order to obtaina broader development and meet harsher service conditions, TiAl-based alloys needto further improve the fracture toughness at room temperature and increase the hightemperature strength. In the present work, toughening and strengthening of theTiAl-based alloy were realized by reinforcing it with continuous Mo fibers. TheMo_f/TiAl composite were successfully fabricated by powder slurry casting andvacuum hot pressing. Effects of the fabrication process parameters on themicrostructure and mechanical properties of the composites were studied; thedeformation and fracture behavior of the composites were observed and analyzed,and the toughening and strengthening mechanism of the composites were obtained.
In this paper, the pre-coupling of the matrix powders and fibers wereconveniently and effectively realized by improved powders slurry casting process.Qualified preforms were prepared by powder slurry with a composition of10gPMMA,70ml acetone and60g Ti and Al powders. The prepared preforms are denseand void-free with uniform fiber distribution and homogeneous dispersion of Ti andAl powders. The composites were eventually obtained by degassing and vacuum hotpressing of the preforms.
In the prepared Mo_f/TiAl composites, the fibers distribute uniformly and haveno obvious change in microstructure after being hot pressed, which maintain themorphology of columnar crystals with high aspect ratio; the matrices mainly containTiAl phase and a small amount of Ti3Al phase. Increasing the hot pressingtemperature and prolonging the hot pressing time can benefit the densification andmicrostructural uniformity of the matrices but exacerbate the interface reaction. There are two continuous reaction layers at the interface of matrix and fibers, whichare δ-(Mo,Ti)3Al phase with columnar grains distributing in a radial way andβ'-(Mo,Ti)A1phase with coarse equiaxed grains, respectively, from the Mo fiber tothe matrix. Nano-indentation test shew that δ phase has higher nano-hardness thanthat of β' phase. The shear strength of the composite interface is higher than367MPa and debonding position is located at the Mo/δ interface. Study on thegrowth dynamics of interface phases indicates that both the interface phases growparabolicly and their grow rates increase with the temperature increasing.
The mechanical properties of the composites are influenced by the fabricationprocess parameters, composite degassed at380℃for1hour and hot pressed at1100℃for1hour has the optimal properties. The composite reinforced byunidirectional fibers has a longitudinal and transverse bending strength of735.4MPa and249.3MPa respectively, and the former is slightly affected by theloading direction. Bending strength of the composite reinforced by orthogonalfibers decreases to374.4MPa. In the temperature range of700-1000℃，bendingstrength of the composites reinforced by unidirectional and orthogonal fibers firstlyincrease and then decrease, and they all reach the maximum value762.9MPa and564.4MPa at800℃.
Compressive yielding strength of the composites decrease with the testingtemperature increasing, and it is slightly higher in the longitudinal direction than inthe transverse direction. The way of fiber distribution has little affect on it.Toughness of the composites reinforced by unidirectional fibers is as high as23.55MPa·m1/2, which is about50%higher than that of the matrix TiAl alloy.
Under the co-effect of the brittle matrix and ductile fibers, the deformation andfracture of the Mof/TiAl composite show different way from that of thehomogeneous materials. The deformation and fracture two courses intertwine withthe change of the load carried by the matrix and fibers. Under room temperaturebending and tension in the longitudinal direction, the deformation and fracturecourses of the composite can be divided into three stages. In the first stage, thedeformation of the whole composite comes from the jointly elastic deformation ofthe matrix and fibers and they carry the load together; in the second stage,deformation of the matrix and fibers mismatch, many cracks initiate and propagatein the matrix in order to match the deformation of the composite, and so the load carried by the matrix decreases, while the fibers continue to be deformed elasticallyand the load carried by them increase; in the third stage, the matrix fails due to thesaturation of the cracks in it, so the fibers carry the load alone and deformation ofthe composite is provided by the elastic and plastic deformation of the fibers, andthe composite fractures after the fracture of the necked fibers. Crack initiates in thematrix and can propagate along the TiAl grain boundaries and across the Ti3Algrains. When meeting the δ and β' two interface phases, cracks can directly goacross them but be prevented by the fibers, the two interface phases deform andcrack together with the matrix. Mo/δ、δ/β' and β'/TiAl interface keep intact beforethe plastic deformation of the fibers, but after that, Mo/δ interface debonds becausethe brittle δ phase and ductile cannot deform jointly. Mo fibers are insensitive tocracks, so the crack in the matrix cannot propagate across the fibers. The fibersfracture when the stress reaches the tensile strength.
At high temperatures, bending deformation and fracture of the composite showdifferent way with that at room temperature. Because of the improvement of theductility and plasticity, the matrix can match the whole deformation of thecomposite not by cracks but by deformation. After the necking and fracture of thefibers, the composite fractures quickly due to the insufficient load carrying abilityof the matrix.
Under longitudinal compressive load, fibers curve greatly under the greatestshear stress with a45°angle to the longitudinal direction without fracture, and thematrix among them breaks with the interface Mo/δ debonding. Under normalcompressive load, the broken surface of the composite due to shear stress is parallelthe fiber axial and has many steps. Under high temperatures, the toughness andplasticity of the composite increase, the composite just deform plastically withoutrupture.
The strengthening and toughening mechanisms have been revealed bycomprehensive analysis of the mechanical properties, deformation and fracturebehavior of the composite. The balance between the high matrix density, goodinterfacial bonding and less fiber damage is a guarantee of the maximum increasingof the composite strength. The toughening mechanisms of the composite include thebridging, plastic deformation, pull-out and debonding of the fibers and thedeflection of the cracks. Among them, the plastic deformation of the fiber makes the Mo fiber have higher toughening effect than that of the brittle ceramic fibers.
本文采用改进的粉末浆料铸造法简便有效地实现了纤维与基体粉末的预复合。选用PMMA、丙酮与Ti、Al混合粉末制备粉末浆料，并在三者配比为10g:70ml:60g时，成功地获得了致密无空洞、纤维排布均匀、Ti、Al粉末分散良好的品质优良的预制体。预制体经除气与热压烧结两步最终制得复合材料。
制备的复合材料中纤维排布均匀，保持了制备前的高长径比的纤维组织形貌；基体以TiAl相为主，并含有少量的Ti3Al相。热压烧结温度升高与时间延长，有利于基体致密化程度及组织均匀性的提高，但加剧界面反应，使界面处生成两个连续的反应层，由Mo纤维到基体依次为呈现放射型柱状晶形貌的δ-(Mo,Ti)3Al相和粗大等轴晶形貌的β'-(Mo,Ti)A1相。纳米压痕测试表明，δ相的硬度及弹性模量高于β'。纤维推出试验测得的界面剪切强度大于367MPa，Mo/δ界面发生脱粘。界面相的生长动力学研究表明，δ相与β'相的生长均遵循抛物线规律，二者的生长速度随温度的升高而加快。
复合材料的性能受制备工艺参数影响，最优工艺参数为380℃保温除气1h，并经1100℃热压烧结1h。制备的单向纤维增强的复合材料的室温纵向与横向弯曲强度分别为735.4MPa与249.3MPa，纵向弯曲强度受加载方向影响很小。正交纤维增强的复合材料的室温弯曲强度下降至374.4MPa。700℃-1000℃范围内，单向纤维与正交纤维增强的复合材料的弯曲强度均随温度的升高先增加后降低，二者均在800℃时达到最高值，分别为762.9MPa与564.4MPa。
复合材料的压缩屈服强度随测试温度的升高而减低；纵向压缩屈服强度比法向略高，纤维排布方式对复合材料的法向压缩强度影响较小。单边缺口梁法测得的单向纤维增强的复合材料的断裂韧性为23.55MPa·m1/2，较基体TiAl提高50%以上。
由于脆性基体与韧性连续纤维的共同作用，Mo_f/TiAl复合材料的变形与断裂过程不同于均质材料，变形与断裂过程相互交织，其间伴随着基体与纤维承载作用的变化。室温纵向弯曲与拉伸时，复合材料的变形断裂过程分为三个阶段：第一阶段，复合材料的整体变形来自于纤维与基体间协同发生的弹性变形，纤维与基体共同承载；第二阶段，基体与纤维变形失配，脆性基体通过裂纹的不断萌生扩展匹配复合材料的变形，其有效承载下降，纤维则继续发生弹性变形，承载不断增加；第三阶段，基体因裂纹数量饱和而完全失效，纤维独自承载，复合材料的变形完全由纤维的弹、塑性变形提供，并最终在纤维颈缩断裂后复合材料完全破坏。复合材料中裂纹萌生于基体，发生沿TiAl相晶界及穿过Ti3Al相晶粒的扩展后，直接穿过δ与β'两界面相而受阻于纤维，两界面相与基体一同变形开裂。Mo/δ、δ/β'与β'/TiAl三界面在纤维塑性变形前均完好无损，之后因脆性δ相与韧性纤维难以协调变形，Mo/δ界面脱粘。Mo纤维对裂纹不敏感，基体裂纹难以扩展穿过纤维，纤维在达到自身的抗拉强度后萌生裂纹并断裂。
高温下复合材料的弯曲变形与断裂方式与室温时不同，塑韧性大为提高的基体能够通过较大的变形而非裂纹的萌生扩展来匹配复合材料的整体变形，在纤维发生颈缩断裂后，复合材料因基体承载能力的不足而快速断裂。
纵向压缩载荷作用下，复合材料中纤维在与纵向成45°方向的最大切应力作用下发生较大弯曲而不断裂，其间基体则沿纤维轴线破碎，并造成界面Mo/δ开裂。法向压缩载荷作用下，复合材料的剪切破坏面平行于纤维轴向，并呈现台阶状。高温时，复合材料塑韧性增加，基体与纤维变形协调一致，复合材料发生墩粗变形而不破裂。
通过对复合材料力学性能及变形与断裂行为的综合分析，揭示了复合材料的强韧化机理：较高的基体致密度、良好的界面结合及较少的纤维损耗三者之间的平衡是复合材料强度得到最大提高的保证；纤维的桥接、塑性变形、脱粘与拔出以及裂纹的偏转枝生是复合材料的主要韧化机制，其中，纤维的塑性变形使强韧性优异的Mo纤维的增韧效果远胜于脆性陶瓷纤维。
High-temperature structural materials play a key role on the realization of therevolutionary development of the aerospace propulsion system. Among them,TiAl-based alloys have become very competitive ones, because they have theadvantages of low density, high strength, good oxidation resistance and excellentcreep resistance. Currently, they have been applied to engine blades, turbine,exhaust valves and other components. However, the rapid development of theaerospace industry proposes higher demands on engine materials. In order to obtaina broader development and meet harsher service conditions, TiAl-based alloys needto further improve the fracture toughness at room temperature and increase the hightemperature strength. In the present work, toughening and strengthening of theTiAl-based alloy were realized by reinforcing it with continuous Mo fibers. TheMo_f/TiAl composite were successfully fabricated by powder slurry casting andvacuum hot pressing. Effects of the fabrication process parameters on themicrostructure and mechanical properties of the composites were studied; thedeformation and fracture behavior of the composites were observed and analyzed,and the toughening and strengthening mechanism of the composites were obtained.
In this paper, the pre-coupling of the matrix powders and fibers wereconveniently and effectively realized by improved powders slurry casting process.Qualified preforms were prepared by powder slurry with a composition of10gPMMA,70ml acetone and60g Ti and Al powders. The prepared preforms are denseand void-free with uniform fiber distribution and homogeneous dispersion of Ti andAl powders. The composites were eventually obtained by degassing and vacuum hotpressing of the preforms.
In the prepared Mo_f/TiAl composites, the fibers distribute uniformly and haveno obvious change in microstructure after being hot pressed, which maintain themorphology of columnar crystals with high aspect ratio; the matrices mainly containTiAl phase and a small amount of Ti3Al phase. Increasing the hot pressingtemperature and prolonging the hot pressing time can benefit the densification andmicrostructural uniformity of the matrices but exacerbate the interface reaction. There are two continuous reaction layers at the interface of matrix and fibers, whichare δ-(Mo,Ti)3Al phase with columnar grains distributing in a radial way andβ'-(Mo,Ti)A1phase with coarse equiaxed grains, respectively, from the Mo fiber tothe matrix. Nano-indentation test shew that δ phase has higher nano-hardness thanthat of β' phase. The shear strength of the composite interface is higher than367MPa and debonding position is located at the Mo/δ interface. Study on thegrowth dynamics of interface phases indicates that both the interface phases growparabolicly and their grow rates increase with the temperature increasing.
The mechanical properties of the composites are influenced by the fabricationprocess parameters, composite degassed at380℃for1hour and hot pressed at1100℃for1hour has the optimal properties. The composite reinforced byunidirectional fibers has a longitudinal and transverse bending strength of735.4MPa and249.3MPa respectively, and the former is slightly affected by theloading direction. Bending strength of the composite reinforced by orthogonalfibers decreases to374.4MPa. In the temperature range of700-1000℃，bendingstrength of the composites reinforced by unidirectional and orthogonal fibers firstlyincrease and then decrease, and they all reach the maximum value762.9MPa and564.4MPa at800℃.
Compressive yielding strength of the composites decrease with the testingtemperature increasing, and it is slightly higher in the longitudinal direction than inthe transverse direction. The way of fiber distribution has little affect on it.Toughness of the composites reinforced by unidirectional fibers is as high as23.55MPa·m1/2, which is about50%higher than that of the matrix TiAl alloy.
Under the co-effect of the brittle matrix and ductile fibers, the deformation andfracture of the Mof/TiAl composite show different way from that of thehomogeneous materials. The deformation and fracture two courses intertwine withthe change of the load carried by the matrix and fibers. Under room temperaturebending and tension in the longitudinal direction, the deformation and fracturecourses of the composite can be divided into three stages. In the first stage, thedeformation of the whole composite comes from the jointly elastic deformation ofthe matrix and fibers and they carry the load together; in the second stage,deformation of the matrix and fibers mismatch, many cracks initiate and propagatein the matrix in order to match the deformation of the composite, and so the load carried by the matrix decreases, while the fibers continue to be deformed elasticallyand the load carried by them increase; in the third stage, the matrix fails due to thesaturation of the cracks in it, so the fibers carry the load alone and deformation ofthe composite is provided by the elastic and plastic deformation of the fibers, andthe composite fractures after the fracture of the necked fibers. Crack initiates in thematrix and can propagate along the TiAl grain boundaries and across the Ti3Algrains. When meeting the δ and β' two interface phases, cracks can directly goacross them but be prevented by the fibers, the two interface phases deform andcrack together with the matrix. Mo/δ、δ/β' and β'/TiAl interface keep intact beforethe plastic deformation of the fibers, but after that, Mo/δ interface debonds becausethe brittle δ phase and ductile cannot deform jointly. Mo fibers are insensitive tocracks, so the crack in the matrix cannot propagate across the fibers. The fibersfracture when the stress reaches the tensile strength.
At high temperatures, bending deformation and fracture of the composite showdifferent way with that at room temperature. Because of the improvement of theductility and plasticity, the matrix can match the whole deformation of thecomposite not by cracks but by deformation. After the necking and fracture of thefibers, the composite fractures quickly due to the insufficient load carrying abilityof the matrix.
Under longitudinal compressive load, fibers curve greatly under the greatestshear stress with a45°angle to the longitudinal direction without fracture, and thematrix among them breaks with the interface Mo/δ debonding. Under normalcompressive load, the broken surface of the composite due to shear stress is parallelthe fiber axial and has many steps. Under high temperatures, the toughness andplasticity of the composite increase, the composite just deform plastically withoutrupture.
The strengthening and toughening mechanisms have been revealed bycomprehensive analysis of the mechanical properties, deformation and fracturebehavior of the composite. The balance between the high matrix density, goodinterfacial bonding and less fiber damage is a guarantee of the maximum increasingof the composite strength. The toughening mechanisms of the composite include thebridging, plastic deformation, pull-out and debonding of the fibers and thedeflection of the cracks. Among them, the plastic deformation of the fiber makes the Mo fiber have higher toughening effect than that of the brittle ceramic fibers.
引文
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