Research progress on thermal load damage behavior of tungsten-based plasma facing materials
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摘要:
在核聚变反应堆运行过程中，面向等离子体钨基材料需要承受住一定次数稳态和瞬态热负荷的冲击而不发生开裂、熔化等损伤，因此改善面向等离子体钨基材料的力学性能及高温稳定性是极其重要的，主要手段包括合金强化、弥散强化、纤维增韧、第二相强化、复合强化等。本文分析了合金强化、弥散强化等改性手段对钨基材料热负荷损伤行为的影响，总结了各种强化手段的优势和不足，并对面向等离子体掺杂钨基材料热负荷损伤行为进行了展望。
Abstract:During the operation of nuclear fusion reactors, the plasma facing tungsten-based materials need to withstand a certain number of steady-state and transient thermal loads without cracking, melting, and other damage. It is very important to improve the mechanical properties and high temperature stability of plasma facing tungsten-based materials, the current modification methods mainly include alloy strengthening, dispersion strengthening, fiber toughening, second phase strengthening, composite strengthening, and so on. The effect of modification methods on the thermal load damage behavior of tungsten-based materials was investigated in this paper, such as alloy strengthening and dispersion strengthening, the advantages and disadvantages of those strengthening methods were summarized, and the thermal load damage behavior of plasma facing tungsten-based materials was prospected.
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Keywords:
- nuclear fusion /
- plasma facing materials /
- thermal load /
- alloying /
- dispersion strengthening
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先进制造业和机器人技术的发展对金属基多孔材料产生了更大的需求和更高的要求。精密机械用含油自润滑轴承、机器人关节摩擦副、电力系统受电滑板等是高性能金属基多孔材料的典型应用领域，不仅要求材料具有良好的减摩耐磨性能，还要具备储油自润滑的功能和良好的电、热传导性能[1-6]。采用粉末烧结技术制备金属基多孔材料具有生产效率高、成分组织易调控、性能稳定等优点，应用前景广阔。
铜基石墨自润滑材料是一种重要的金属基多孔材料，不仅具有铜基合金良好的力学性能、耐蚀性能、导热和导电性能，还具备减摩、耐磨、抗咬合等摩擦学性能，在航空、汽车、铁路、电力、矿山等领域中获得了广泛的关注[4-6]。由于铜与石墨既不润湿也不发生反应，所以单纯的铜–石墨复合材料界面结合强度低，在承受较重载荷时，易造成石墨相剥落，使摩擦系数增大，磨损率增加[7]。为了改善铜基体与石墨的润湿性，提高铜–石墨复合材料的综合力学性能，人们采取了铜基体合金化、石墨颗粒表面处理、氧化物掺杂等措施[7-10]。然而，铜基体合金化会降低材料导电性和耐热稳定性，石墨颗粒表面处理会增加工艺难度和成本。因此，加入不溶于铜基体的表面活性组分就成为改善铜–石墨界面润湿性的新途径[11]。
SnO2与石墨有着良好的亲和性，二者的复合物可用于气敏传感器、太阳能电池、薄膜电阻、光催化触媒和锂离子电池负极等材料[12-14]。此外，SnO2在高温下可以被碳还原为液态锡，液态锡对铜有着非常好的浸润性，可以溶入铜中形成固溶体。因此，将SnO2加入铜–石墨复合材料有助于改善石墨与铜的界面结合。另外，SnO2是一种具有导电性的氧化物，将其加入铜–石墨复合材料不会对材料的导电性造成太大的损失。基于上述考虑，本文以Cu、石墨和SnO2粉末为原料，采用SiO2–B2O3–Al2O3系铜合金助焊剂辅助的常压粉末烧结工艺制备了Cu–C–SnO2多孔复合材料，前期已报道了SiO2–B2O3–Al2O3系助焊剂对Cu–C–SnO2多孔复合材料组织和性能的影响[15]，本文探讨烧结温度对其组织和性能的影响。
1. 实验材料及方法
1.1 实验材料
实验中使用的原料为400目雾化铜粉（纯度≥99.7%）、高纯石墨粉（纯度≥99.9%）、氧化锡粉（纯度≥99.0%）、SiO2–B2O3–Al2O3系铜合金助焊剂（融化温度约770℃）。为了保证烧结体中有足够量的石墨不被二氧化锡氧化烧损殆尽，又不至于因石墨太多而不利于烧结，依据经验取C/SnO2质量比为1.0:6.3；为了保证在烧结体中铜颗粒连成骨架的情况下掺入尽可能多的SnO2和石墨，参考文献[16]，取Cu/(Cu+C+SnO2)体积分数为0.67。以Cu密度8.96、石墨密度2.26和SnO2密度6.48计算，按照Cu/C/SnO2质量比4.091:0.159:1.000进行粉体配料。用JA1003型千分之一电子天平进行配料称重。先将SnO2和石墨粉在研磨钵中混合、研磨10 min，再加入铜粉，研磨20 min；然后，按照以往研究确定的助焊剂最佳加入量（混合粉体与助焊剂质量比100:5[15]）加入一定质量的SiO2–B2O3–Al2O3系助焊剂，继续研磨20 min，得到成品粉体。
1.2 烧结工艺
将配制好的成品粉体填充于石墨型中，橦实并封盖，在SX2-2.5-10A型箱式电阻炉中加热、保温和随炉冷却，烧结工艺如图 1所示。
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图 1 Cu–C–SnO2混合粉体烧结工艺：（a）烧结粉体封装；（b）烧结温度–时间曲线Figure 1. Sintering technology of the Cu–C–SnO2 mixture powders: (a) the capsulation of the mixture powders; (b) the sintering temperature-time curve1.3 性能测试与表征
把烧结好的样品从石墨型中取出，将其两个端面磨平抛光，然后进行组织观察和性能测试。金相组织观察、X射线衍射分析、烧结线收缩率测定、空气粘性渗透系数测试、电阻率测试、硬度测试等采用的仪器和方法与作者以往研究[15]所采用的相同。样品密度通过几何尺寸测量和质量称重数据计算得出，如式（1）所示。
$$ \rho = \frac{{12M}}{{{\rm{ \mathsf{ π} }}h\left( {{D^2} + {d^2} + D \cdot d} \right)}} $$ (1) 式中：ρ为样品密度，g·cm-3; h为样品高度，cm; D为样品大端直径，cm; d为样品小端直径，cm; M为样品质量，g。
样品的渗油率测试装置和方法如图 2所示。试样由弹簧钢丝夹和纤维丝悬挂于油液上方，调节张丝旋钮，使试样下端面与油液表面恰好接触。油液表面积足够大，以使油液表面高度在渗油实验过程中的变化可以忽略。用称重电子天平记录试样重量随渗油时间的变化曲线，直至试样渗油达到饱和（重量不再增加），称量试样渗油饱和后的质量，按式（2）计算试样的渗油率。
$$ k = \frac{{{M_{\rm{o}}} - M}}{M}\rho $$ (2) 式中：k为试样的渗油率，g·cm-3; M为试样的原始质量，g; Mo为试样渗油达到饱和时的质量，g；ρ为试样的原始密度，g·cm-3。实验用油的动力粘度为0.065 Pa·s，密度为0.925 g·cm-3。
2. 结果与讨论
2.1 烧结温度对烧结体金相组织的影响
烧结体经磨平抛光后直接在光学金相显微镜下进行观察，结果如图 3所示，其金相组织主要由三种不同的物相组成：明亮的金属铜相（Cu）、灰暗的氧化物陶瓷相、黑色的石墨相（C）或孔洞[15]。可以看出，随着烧结温度的升高，烧结体中铜相颗粒之间界面逐渐减少，颗粒之间相互融合，连成一体；尤其是烧结温度从900℃升高到950℃时，铜相颗粒发生显著的再结晶，晶粒尺寸剧烈长大，孔洞轮廓变得光滑圆整。Felege等[17]最新的研究工作发现，将主要粒径30μm的球形纯铜粉末经300 MPa单轴压力压实，在氢气保护下进行烧结实验，在1020℃烧结温度下铜颗粒发生了明显的再结晶，Σ3孪晶和大角度晶界明显减少，孔洞形状球化，大小趋于均匀。
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图 3 不同烧结温度下烧结体金相组织：（a）750℃；（b）800℃；（c）850℃；（d）900℃；（e）950℃Figure 3. Metallographic structures of the sintered compacts at different sintering temperatures: (a) 750℃; (b) 800℃; (c) 850℃; (d)900℃; (e) 950℃图 4为不同烧结温度下所得烧结体的X射线衍射分析结果。由图 4（a）可以看出，烧结体中氧化物相的成分和结构比较复杂，有Sn O、SnO2、Cu2O、SiO2、莫来石和刚玉。烧结温度为750℃时，几乎没有SiO2、莫来石和刚玉等物相的衍射峰出现；随着烧结温度升高，Sn O、SnO2、Cu2O逐渐减少，SiO2、莫来石和刚玉等物相的衍射峰逐渐增强，说明硅、铝氧化物矿化程度增大；当烧结温度达到950℃时，Sn O、SnO2、Cu2O基本消失。图 4（b）为各相质量分数随烧结温度的变化情况，由X射线衍射主峰强度，以Cu相为内标计算所得。
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图 4 烧结体X射线衍射分析：（a）衍射谱；（b）物相含量Figure 4. X-ray diffraction analysis of the sintered compacts: (a) diffraction spectrum; (b) phase contentSn O、SnO2、Cu2O随烧结温度升高而减少的原因是发生了如下式（3）~式（4）的碳还原反应。
$$ {{\rm{SnO}} + {\rm{C}} = {\rm{Sn}} + {\rm{CO}}} $$ (3) $$ {{\rm{Sn}}{{\rm{O}}_2} + {\rm{C}} = {\rm{SnO}} + {\rm{CO}}} $$ (4) $$ {{\rm{C}}{{\rm{u}}_2}{\rm{O}} + {\rm{C}} = 2{\rm{Cu}} + {\rm{CO}}} $$ (5) 根据热力学计算[18]，在标准状态下，式（3）和式（4）在700℃以上就可以发生，式（5）在100℃以上就可以发生。在本研究所采用的烧结条件下，CO的平衡分压力远小于1标准大气压，上述反应的热力学条件充分。但由于反应物都是固相，需要较高的温度提供反应激活能，因此，在有限的烧结时间里，上述反应进行的程度随烧结温度的变化而变化，烧结温度越高，反应进行的越充分，氧化物残留量越少。
硅、铝氧化物矿化是不同氧化物固体颗粒在低熔点熔融复合盐的包围下通过界面反应进行的，需要反应界面两侧氧化物的不断溶解和扩散。氧化物溶解和扩散的速度随温度升高而加快，在有限的烧结时间里，氧化物溶解和扩散进行的程度随烧结温度的升高而增大，所生成的矿化晶体（SiO2、莫来石和刚玉等）数量随烧结温度的升高而增加。
2.2 烧结温度对烧结体线收缩率和密度的影响
图 5和图 6分别是烧结体线收缩率和密度随烧结温度的变化情况，两者趋势相同。随烧结温度升高，烧结体线收缩率和密度增大。烧结温度800℃为一临界点，烧结温度低于800℃时，线收缩率和密度随烧结温度升高迅速增大；烧结温度高于800℃时，线收缩率和密度随烧结温度的升高而增加速率降低，并呈指数增加趋势。出现上述变化的原因是存在着两种烧结收缩机制：800℃以下烧结时，由助焊剂熔化造成的氧化物熔融凝固收缩是导致烧结体收缩的主要因素，烧结温度升高，氧化物熔融程度增加，冷却时凝固收缩量增大；800℃以上烧结时，由助焊剂熔化造成的氧化物熔融已经彻底完成，继续升高烧结温度，氧化物熔融程度不再增加，冷却时凝固收缩量也就不再增加，这时铜颗粒固相再结晶造成的收缩表现出来。固相再结晶是减少晶体缺陷而产生收缩，其收缩量远小于凝固收缩量。固相再结晶通过晶体内的原子扩散来实现，需要热激活，其进行的速度受温度控制，与温度成指数关系，在有限的烧结时间里，铜颗粒固相再结晶进行的程度也就与烧结温度成指数关系，由此造成烧结体收缩量随烧结温度升高呈指数增加的趋势。
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图 5 烧结温度对烧结体收缩率的影响Figure 5. Effect of sintering temperature on the sintering shrinkage of the sintered compacts![]()
图 6 烧结温度对烧结体密度的影响Figure 6. Effect of sintering temperature on the density of the sintered compacts2.3 烧结温度对烧结体渗透性和导电性的影响
图 7、图 8和图 9分别为烧结体电阻率、渗油率以及空气粘性渗透系数随烧结温度的变化情况。从图中可以看出，随着烧结温度升高，烧结体的电阻率、渗油率、空气粘性渗透系数呈现出相似的变化趋势：先减小后增大，在烧结温度850~900℃之间达到最小值。
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图 7 烧结温度对烧结体电阻率的影响Figure 7. Effect of sintering temperature on the electrical resistivity of the sintered compacts![]()
图 8 烧结温度对烧结体渗油率的影响Figure 8. Effect of sintering temperature on the oil penetration rate of the sintered compacts![]()
图 9 烧结温度对烧结体空气粘性渗透系数的影响Figure 9. Effect of sintering temperature on the air permeability coefficient of the sintered compacts渗油率和空气粘性渗透系数两者都反映了烧结体通透孔隙率的大小，理论上二者都与通透孔隙率成正比，但由于油和空气对烧结体物相润湿性存在显著差异，在孔隙尺寸很小时，毛细力作用对渗油率和空气粘性渗透系数的影响明显不同。油对于烧结体物相是润湿的，毛细力的作用促进油的渗透，烧结体孔隙尺寸越小，油越容易渗透；空气粘性渗透的阻力却随烧结体孔隙尺寸减小而增大。当烧结温度升高到较高范围（900~950℃）时，烧结体收缩使孔隙尺寸减小，SnO2和石墨相烧蚀使分散细小的孔隙数量增加，从而使烧结体渗油率显著回升，而空气粘性渗透系数却回升缓慢。电阻率随烧结温度变化的趋势与渗油率和空气粘性渗透系数的变化趋势相似，说明影响烧结体导电性的主要因素是材料的孔隙率。
2.4 烧结温度对烧结体硬度的影响
图 10为烧结体里氏硬度随烧结温度的变化。可以看出，随着烧结温度升高，烧结体里氏硬度经历了两个上升阶段，在两个上升阶段之间存在一个硬度突变。在750~800℃烧结温度范围内，里氏硬度随烧结温度升高而升高；在烧结温度800~850℃之间存在一个临界温度，里氏硬度突然降低；在烧结温度850~950℃范围内，里氏硬度继续随烧结温度升高而升高。分析其原因，里氏硬度表征的是材料抵抗冲击性动载荷的能力，对于易变形的金属和分散的陶瓷颗粒堆积体示值都低，对于坚硬强韧的致密物相示值就高。虽然热力学计算表明式（5）在本研究所采用的烧结条件下可以自发进行，但动力学上Cu2O的生成与分解存在一个临界温度，图 11所示的金属铜在大气中的微商热重曲线（derivative thermogravimetry, DTG）表明此温度为827℃。温度高于此临界温度，Cu和Cu2O随气氛氧势不同而相互转化的速度很快；温度低于此临界温度，Cu和Cu2O随气氛氧势不同而相互转化的速度很慢。在本研究所采用的烧结气氛（有碳存在）下，烧结温度低于827℃时，Cu2O可以稳定存在，随烧结温度升高，Cu氧化程度增大，Cu2O数量增加（见图 4），烧结体的里氏硬度随之升高；烧结温度高于827℃时，Cu2O开始被C还原成金属Cu，烧结体里氏硬度突然降低，随着烧结温度继续升高，Cu颗粒固相再结晶使烧结体组织致密化，再加上硅、铝氧化物矿化数量增加，从而使里氏硬度随之继续升高。
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图 10 烧结温度对烧结体硬度的影响Figure 10. Effect of sintering temperature on the hardness of the sintered compacts![]()
图 11 金属铜在大气中的微商热重曲线Figure 11. Derivative thermogravimetry curve of copper in the atmosphere3. 结论
（1）随着烧结温度升高，烧结体中的铜相颗粒因再结晶而逐渐融合，颗粒之间界面减少；尤其是烧结温度从900℃升高到950℃时，铜相颗粒发生显著的再结晶。
（2）随着烧结温度升高，Sn O、SnO2、Cu2O逐渐减少，硅、铝氧化物矿化程度增大；烧结温度达到950℃时，Sn O、SnO2、Cu2O基本消失。
（3）随着烧结温度升高，烧结体线收缩率和密度增大。烧结温度800℃为一临界点，烧结温度低于800℃时，线收缩率和密度随烧结温度升高迅速增大；烧结温度高于800℃时，线收缩率和密度随烧结温度升高而呈指数趋势缓慢增加。
（4）随着烧结温度升高，烧结体的电阻率、渗油率、空气粘性渗透系数呈现出相似的变化趋势：先减小后增大，在烧结温度850~900℃之间达到最小。
（5）随着烧结温度升高，烧结体里氏硬度经历了两个上升阶段，在烧结温度800~850℃之间存在一个临界温度，里氏硬度突然降低。
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图 4 不同吸收功率密度下，单脉冲5 ms后S–WZC和R–WZC加载区域的扫描电子显微形貌[53]：（a）0.44 GW·m−2；（b）0.66 GW·m−2；（c）和（d）0.88 GW·m−2；（e）和（f）1.10 GW·m−2
Figure 4. SEM images of S–WZC and R–WZC in the loaded areas after single pulse for 5 ms at room temperature with the various absorbed power densities[53]: (a) 0.44 GW·m−2; (b) 0.66 GW·m−2; (c) and (d) 0.88 GW·m−2; (e) and (f) 1.10 GW·m−2
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图 7 100次类边界局域模式脉冲后热压和高能率锻造W–Y2O3试样表面形貌[68]：（a）热压，0.22 GW·m−2；（b）热压，0.33 GW·m−2；（c）高能率锻造，0.55 GW·m−2；（d）高能率锻造，0.66 GW·m−2
Figure 7. Surface morphology of the W–Y2O3 samples by hot pressing (HP) and HERF after 100 pulses[68]: (a) HP, 0.22 GW·m−2; (b) HP, 0.33 GW·m−2; (c) HERF, 0.55 GW·m−2; (d) HERF, 0.66 GW·m−2
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