Research progress on C-coated Ti composite powders used for preparing high-performance Ti matrix composites
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摘要:
钛基复合材料中增强相的形貌和分布是决定材料性能的关键，常规粉体机械混合后烧结引入增强相的方式存在形貌难调控、分布单一且均匀性差等问题，导致其强化效果不佳。针对该问题，本团队开发了一系列碳包覆钛复合粉体，通过设计包覆碳源的结构与组成调控粉体烧结过程中增强相的形成路径，不仅实现了增强相形貌调控和不同形貌的组合搭配，而且得到了晶内和晶界双增强相组织，大幅提升了钛基复合材料的力学性能。在此基础上，将碳包覆钛复合粉体拓展应用至钛基复合材料的3D打印领域，解决了高品质复合粉体缺乏并制约其发展的瓶颈问题。总结并评述了碳包覆钛复合粉体在制备钛基复合材料中取得的研究结果与工作进展，为增强相设计与调控提供新的研究思路及技术路线。
Abstract:The morphology and distribution of reinforcements in the titanium matrix composites (TMCs) are crucial in determining the material performances. Due to the uncontrollable morphology and inhomogeneous distribution of the reinforcements in the current TMCs, a series of C-coated Ti composite powders were developed by fluidization technology. By designing the structure and composition of C-coatings, the different morphology combinations and the intragranular/interface reinforcements were both achieved, which significantly improved the mechanical properties of TMCs. Furthermore, the composite powders were also extended to the 3D printing of TMCs, which solved the bottleneck issue of lacking the high-quality composite powders. The research results and work progress of the C-coated Ti composite powders in the preparation of high-performance TMCs were summarized and reviewed, providing the new insight and technical route for the design and control of reinforcements in TMCs.
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碳化钛（TiC）具有高硬度、高熔点、导电性好、耐腐蚀、抗高温等优点，被广泛应用于工业工程、航空航天、核工业等领域[1−3]。由于TiC具有强的共价键，烧结性较差，影响了TiC陶瓷的力学性能，并限制了TiC陶瓷的应用[4]。通常，通过添加第二相（WC、ZrC、SiC、TiN等）以及金属相（Ti、Mo、Co、Ni、Cr等）改善TiC陶瓷的烧结性，提高陶瓷力学性能[5−7]。氮化钛（TiN）具有硬度高、熔点高、化学稳定性好、摩擦系数低、导电性能好、颜色独特且可变等特点，被广泛应用于机械工业、生物医疗、导电材料等领域[8]。在一定条件下，TiN可与TiC形成TiCxNy固溶体，TiCxNy固溶体的韧性和化学稳定性优于TiC，硬度和耐磨性优于TiN，故将两者复合形成固溶体，可兼容TiC和TiN的优势[9−10]。此外，WC、HfN与TiC、TiN或TiCN有较好的物理化学相容性[11−12]，它们是TiC、TiN或TiCN陶瓷材料的理想增强相。金属相不仅可改善TiC、TiN陶瓷材料的微观组织，还可提高材料的力学性能。金属Ni对TiC和TiB2陶瓷材料有较好的润湿性，随着Ni含量的增加，TiC–TiB2材料的硬度、抗弯强度和断裂韧度均有所提高[13]；适量的Mo能够细化TiC–TiN–WC陶瓷材料的晶粒且能提高陶瓷的抗弯强度[9]；Ti作为TiC陶瓷材料的粘结剂，可以使材料获得高的相对密度[14]。金属Re是一种熔点高、稳定性好的金属，也是陶瓷材料的理想添加剂。Zi等[15]发现Re可改善Ni与Al2O3陶瓷间的润湿性。Marcin和Anna[16]发现在Cr–Al2O3复合材料中加入Re可提高材料的摩擦磨损性能。但是，目前有关Re对TiC、TiN、TiCN陶瓷材料性能影响方面的报道较少。
本研究以TiC和TiN为基体，以WC和HfN为增强相，以Ni和Re为金属相，通过热压烧结技术制备TiCN–WC–HfN陶瓷，研究Re含量（摩尔分数）对材料微观组织和力学性能的影响。
1. 实验材料及方法
1.1 实验材料及制备
制备TiCN–WC–HfN陶瓷所用TiC、TiN、WC、HfN、Ni、Re粉末均来自上海允复纳米科技有限公司，粉末平均粒径均为0.5 μm，纯度均大于99%，具体组分及含量见表1。
表 1 TiCN–WC–HfN陶瓷组分及含量（摩尔分数）Table 1. Composition and content of the TiCN–WC–HfN ceramics% 材料编号 TiC TiN WC HfN Ni Re R0 30 30 15 15 10.0 0 R1 30 30 15 15 8.0 2.0 R2 30 30 15 15 7.5 2.5 R3 30 30 15 15 7.0 3.0 根据表1称量原料粉末，置于球磨罐中进行球磨，球磨介质为硬质合金球和无水乙醇，球磨时间72 h。将球磨后的浆料置入干燥箱中干燥。随后，用100目的网筛过筛，倒入直径为50 mm的石墨模具中，完成素坯的制备。使用ZT-40-20型真空热压烧结炉烧结素坯，其中烧结温度为1550 ℃，保温时间为60 min，升温速率为10 ℃·min−1，烧结压力为30 MPa。烧结后的材料经切割、粗磨、细磨、抛光等工艺制成3 mm×4 mm×40 mm的试样条。
1.2 表征方法
依据GB/T6569-2006[17]采用三点抗弯法在CREE-8003G材料试验机上测试材料的抗弯强度，其跨距为30 mm，加载速度为0.5 mm·min−1。依据GB/T16534-2009[18]在HVS-30硬度计上测试材料的维氏硬度，加载载荷196 N，保压15 s。采用压痕法[19]测试材料的断裂韧度。力学性能的测试均以15个测试结果的算术平均值作为测试结果。使用RAY-10AX-X-ray型X射线衍射仪（X-ray diffraction，XRD）和能谱仪（energy disperse spectroscope，EDS）分析材料的物相组成，并通过Supra-55型扫描电镜（scanning electron microscope，SEM）观察材料的抛光面和断口形貌。
2. 结果与讨论
2.1 物相组成
图1是所制备TiCN–WC–HfN（R3）陶瓷的X射线衍射图谱。由图1可见，陶瓷试样的主要相为TiC0.41N0.50、WC、HfN和TiC，同时，含有少量的Ni和Re。图谱中未发现TiN，而有大量TiC0.41N0.50固溶体和一定量TiC，这表明在热压烧结过程中，几乎所有TiN与大部分TiC发生了固溶，形成了TiC0.41N0.50固溶体。Verma等[20]在研究TiCN基陶瓷材料时发现了少量(Ti,W)(C,N)固溶体，但在本研究的X射线衍射图谱中并未发现，可能是其含量较少，无法被检测到。
图2是TiCN–WC–HfN陶瓷的抛光面形貌及相组成。由图2（a）可见，材料由黑色相、白色相、浅灰色相和深灰色相组成。由图2（b）可见，深灰色相所占面积最大，其次依次为浅灰色、白色相和黑色相。图3是各相的能谱分析。由图3（a）可见，黑色相中的C和Ti元素含量较高，其中C的原子数分数为42.88%，Ti的原子数分数为42.40%，其比接近1:1，结合X射线衍射图谱分析可知，黑色相的主要成分是TiC。同理，由图3（b）可见，白色相中N和Hf元素含量较高，其中N的原子数分数为43.09%，Hf的原子数分数为41.52%，其比接近1:1，结合X射线衍射图谱结果可知，白色相的主要成分是HfN。由图3（c）可见，浅灰色相中C和W元素含量较高，其原子数分数分别为47.31.%和44.92%，其比接近1:1，结合X射线衍射图谱分析可知，浅灰色相的主要成分是WC。浅灰色相的边界较为平直，晶粒形貌近似四边形，这与杨方等[21]所报道的WC形貌基本一致。由图3（d）可见，深灰色相中Ti、C和N原子含量较高，其原子数分数分别为48.41%、19.27%和25.06%，其比接近1.0:0.4:0.5，结合X射线衍射图谱分析可知，深灰色相的主要成分是TiC0.41N0.50。此外，陶瓷相在液相金属Ni和Re中完成溶解–析出–结晶后，Ni和Re会粘附在晶粒周围；同时，在烧结压力的作用下，液相金属Ni和Re会填充到晶粒间的空隙中；Ni和Re在黑色相、白色相、浅灰色相和深灰色相的能谱中均有体现，但其含量相对较低。
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图 2 TiCN–WC–HfN陶瓷显微形貌（a）及相组成（b）Figure 2. Microstructure (a) and phase composition (b) of the TiCN–WC–HfN ceramics![]()
图 3 TiCN–WC–HfN陶瓷各相能谱分析:（a）黑色相；（b）白色相；（c）浅灰色相；（d）深灰色相Figure 3. EDS analysis of the TiCN–WC–HfN ceramics: (a) black phase; (b) white phase; (c) light gray phase; (d) gray phase2.2 TiCN–WC–HfN陶瓷的微观组织
图4是TiCN–WC–HfN陶瓷的断口形貌。由图可见，随着Re摩尔分数从0增到3.0%，晶粒呈先变大后变小的趋势，甚至发生了晶粒聚集，如图中虚线框所示，这表明Re在一定程度上具有抑制晶粒长大的作用。同时，在试样R0～试样R3中均存在凹坑，如图中实线圆圈所示，且试样R2中的凹坑最多。这些凹坑是由材料中小晶粒拔出所致（如图中箭头所示）；在材料的断裂过程中，这些小晶粒可起到钉扎作用，这有利于材料力学性能的提高。另外，试样中均存在解理面（如图中实线框所示），试样R0中的解理面相对较少，晶粒断面相对平整；而试样R1～试样R3中的解理面较多，这表明晶粒断裂时非一次性直接断裂，而是在外力的作用下逐渐断裂，这种断裂会消耗更多的断裂能，有利于材料抗弯强度和断裂韧度的提高。试样R0和试样R3存在晶粒聚集现象，其中试样R3中的晶粒发生了严重聚集，这会削弱材料的力学性能。
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图 4 TiCN–WC–HfN陶瓷断口形貌：（a）R0；（b）R1；（c）R2；（d）R3Figure 4. Fracture morphologies of the TiCN–WC–HfN ceramics: (a) R0; (b) R1; (c) R2; (d) R32.3 TiCN–WC–HfN陶瓷的力学性能
图5是Re含量（摩尔分数）对TiCN–WC–HfN陶瓷力学性能的影响。由图可见，当Re的摩尔分数由0增到3.0%时，材料的硬度、抗弯强度和断裂韧度均先增大后减小；当Re的摩尔分数为2.5%时，材料的力学性能最优，其维氏硬度为(19.25±0.21) GPa、抗弯强度为(1304±23) MPa、断裂韧度为(7.73±0.22) MPa∙m1/2；而当Re摩尔分数为0时，材料的力学性能分别为(18.04±0.18) GPa、(1021±19) MPa和(7.11±0.19) MPa∙m1/2。当Re摩尔分数为2.5%时，材料在断裂过程中，其断口上较多的小晶粒被拔出形成凹坑，以及晶粒在断裂过程中形成解离面都需要消耗大量的断裂能，这是其力学性能较高的主要原因。当Re摩尔分数为3.0%时，材料的抗弯强度和维氏硬度发生了较大幅度的降低，这是由晶粒的严重聚集造成的。
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图 5 Re含量对TiCN–WC–HfN陶瓷力学性能的影响Figure 5. Relationship between the Re content and mechanical properties of the TiCN–WC–HfN ceramics图6为TiCN–WC–HfN（R3）陶瓷的裂纹扩展路径。由图可见，裂纹扩展时发生了偏转和桥连。裂纹偏转和裂纹桥连会消耗更多的断裂能，这有利于材料断裂韧度的提高[22]。此外，在裂纹扩展时，存在沿晶扩展和穿晶扩展，即材料在断裂时发生了沿晶断裂和穿晶断裂，这种沿晶与穿晶并存的断裂方式也有助于材料断裂韧度的提高[23]。
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图 6 TiCN–WC–HfN（R3）陶瓷裂纹扩展路径Figure 6. Crack propagation of the TiCN–WC–HfN (R3) ceramics3. 结论
（1）烧结后的TiCN–WC–HfN–Ni–Re陶瓷材料由TiC0.41N0.50、WC、HfN、TiC、Ni和Re组成，其中TiC0.41N0.50是TiC与TiN在烧结过程中生成的固溶体。
（2）在TiCN–WC–HfN陶瓷的断口上存在凹坑和解理面。当Re摩尔分数为0时，其断口上的解理面相对较少，晶粒的断面相对平整；当Re摩尔分数为2.5%时，材料断口上的凹坑较多；当Re摩尔分数为0和3.0%时，材料断口上存在晶粒聚集的现象。
（3）当Re摩尔分数由0增到3.0%时，材料的硬度、抗弯强度和断裂韧度均先增大后减小。当Re摩尔分数为2.5%时，材料的力学性能最优，其维氏硬度为(19.25±0.21) GPa、抗弯强度为(1304±23) MPa、断裂韧度为(7.73±0.22) MPa∙m1/2。材料在断裂过程中存在穿晶断裂和沿晶断裂，裂纹发生了偏转和桥连。
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图 1 不同类型碳包覆钛复合粉体的制备方法示意图
Figure 1. Schematic diagram of the fabrication process for the different C-coated Ti powders
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图 2 不同碳包覆钛粉体显微形貌[29‒32]：（a）A–C/TC4粉的表面微观形貌；（b）A–C/TC4粉的截面微观形貌；（c）CNTs/Ti粉的微观形貌；（d）CNTs的透射电镜图片；（e）H–C/TC4粉的微观形貌；（f）两种碳源的透射电镜图片
Figure 2. Morphologies of the different C-coated Ti powders[29‒32]: (a) SEM image of the A–C/TC4 powder surface; (b) sectional SEM image of the A–C/TC4 powder; (c) SEM image of CNTs/Ti powder; (d) TEM image of the extracted CNTs; (e) SEM image of H–C/TC4 powder; (f) TEM image of the extracted CNTs and A–C
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图 3 A‒C/TC4粉的热压烧结组织及增强相[29]：（a）和（b）A–C/TC4粉体的烧结组织和增强相分布的扫描电镜照片；（c）和（d）α-Ti晶内分布纳米片状增强相的形貌及物相分析；（e）和（f）β-Ti晶内分布颗粒增强相的形貌及物相分析
Figure 3. Sintered A–C/TC4 samples and the reinforcements[29]: (a) and (b) SEM images of the sintered A–C/TC4 samples and the reinforcement distribution; (c) and (d) the morphology and phase analysis of the nanoplatelets inside α-Ti grains; (e) and (f) the morphology and phase analysis of the nanoparticles inside β-Ti grains
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图 4 CNTs/Ti粉在不同温度的烧结组织和增强相形貌分布[31]：（a）和（b）900 ℃；（c）和（d）1000 ℃；（e）和（f）900 ℃烧结样品中保留的CNTs和形成的TiC颗粒
Figure 4. Microstructure of the CNTs/Ti powder samples sintered at different temperatures[31]: (a) and (b) 900 ℃; (c) and (d) 1000 ℃; (e) and (f) TEM images of the remained CNTs and the formed TiC nanoparticles in the CNTs/Ti powder samples sintered at 900 ℃
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图 6 CNTs/TC4粉体的热压烧结组织
Figure 6. Microstructures of the samples fabricated from the CNTs/TC4 powders by hot press sintering
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图 7 H‒C/TC4粉体烧结组织及增强相[32]：（a）和（b）H–C/TC4粉体烧结组织及增强相分布；（c）～（e）晶界增强相的透射电子显微镜照片；（f）晶内增强相的透射电子显微镜照片
Figure 7. Sintered H–C/TC4 samples and the reinforcements[32]: (a) and (b) sintered H–C/TC4 sample microstructures and the reinforcement distribution; (c)～(e) TEM images of the interfacial reinforcements; (f) TEM images of the intragranular reinforcements
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图 8 不同形貌和分布组合增强相强化钛基复合材料[29‒32]：（a）与文献报道中钛基复合材料压缩屈服强度对比；（b）晶界增强相的强化机制；（c）晶内增强相的强化机制；（d）晶内/晶界双增强相组织的强化机制示意图
Figure 8. Reinforcements with the different morphologies and distribution combinations in TMCs[29‒32]: (a) comparison of the TMCs compressive yield strength reported in literatures; (b) strengthening mechanisms of the interfacial reinforcements; (c) strengthening mechanisms of the intragranular reinforcements; (d) strengthening mechanism of the interfacial/intragranular double reinforced phase
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图 9 CNTs/GA–TC4粉体的扫描电镜形貌（a）、原位合成CNTs扫描电镜形貌（b）、复合粉体打印样品宏观形貌（c）及打印样品的三维重构照片（d）[47]
Figure 9. SEM image of the CNTs/GA–TC4 powders (a), SEM image of the in situ synthesized CNTs (b), macro-profile of the samples printed by the CNTs/GA–TC4 powders (c), and 3D reconstructed images of the printed sample (d)[48]
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