铅冷快堆是最具应用前景的四代堆型之一。Pb-Bi共晶(lead-bismuth eutectic，LBE)由于熔点低(125℃)、沸点高(1670℃)、良好的中子经济性和化学惰性等，是铅冷快堆的首选冷却剂。铅冷快堆的设计服役温度高达550℃，结构材料面临耦合快中子辐照、液态LBE腐蚀、液态金属脆化(LME)等损伤问题，特别要求堆芯结构材料能够承受50~150 dpa (displacements per atom)辐照损伤、具有良好的长时热老化力学性能(拉伸强度、塑性、断裂韧性、蠕变与疲劳等)和良好的液态LBE环境相容性、低活化性等[1~3]，因此结构材料成为制约铅冷快堆研发与应用的瓶颈。最近Gong等[3]系统综述了铅冷快堆的候选结构材料及其液态LBE环境相容性，指出铁素体/马氏体钢(T91、HT9等)、奥氏体不锈钢(316、15-15Ti等)、氧化物弥散强化钢、MAX相陶瓷与碳化硅等是主要的候选材料，其中铁素体/马氏体钢与奥氏体不锈钢是核反应堆已大量应用的结构材料，积累了大量的服役性能数据，并收录在美国机械工程师协会(ASME)等相关手册中，是近期最有可能应用于铅冷快堆的结构材料。因此，铁素体/马氏体钢与奥氏体不锈钢的液态LBE环境相容性成为国内外研究的重点。研究[3~12]表明，T91与316钢在低于450℃、含氧的液态LBE环境中，表面可生成保护性氧化膜，具有良好的耐腐蚀性能；但当温度超过450℃时，表面氧化膜疏松多孔，且容易发生剥落，特别是溶解氧(DO)溶度低于10-8% (质量分数，下同)时会发生严重的溶解腐蚀，服役性能急剧恶化。同时，铁素体/马氏体钢在150~450℃液态LBE环境中，可能发生严重的LME，导致塑性、断裂韧性、蠕变强度和疲劳强度等显著降低[3,13~25]；尽管316钢在液态LBE环境中的LME敏感性低，但其在高温低氧条件下，Ni元素发生选择性溶解导致基体转变成铁素体[3,8,12,13]，可能降低其力学性能。因此，T91与316钢如果作为铅冷快堆的结构材料，需提高其与液态LBE的相容性。优化控制液态LBE中的DO浓度是改善材料液态LBE腐蚀性能的一种有效途径。根据Ellingham图，液态LBE中的DO浓度的选择，既要保证铁基结构材料表面能够生成Fe3O4，又要确保不生成PbO，一般应控制在10-8%~10-6%[1~3]。在此浓度范围内，需综合考虑DO浓度对氧化膜生长热力学与动力学的影响，如何确定铅冷快堆服役时的最优DO浓度，尚需要大量材料环境相容性数据来支撑。合金化也是提高结构材料耐LBE腐蚀性能的有效途径。通过添加Si、Al等元素，生成富Si或Al的氧化物，提高了氧化膜的致密性及其与基体的结合力，成为改善铁素体/马氏体钢和不锈钢的耐液态LBE腐蚀性能的主流方式[26~35]。然而，Si添加会促进有害相的形成，损伤合金的力学性能。Chen和Rong[36]研究表明，Si含量增加导致铁素体/马氏体钢中Laves相析出，提高了韧脆转变温度。Van Den Bosch等[37]研究表明，Si添加会增加铁素体/马氏体钢的LME敏感性。Gong等[38]结合实验与第一性原理计算，揭示了Si添加导致bcc结构纯Fe LME敏感性增加的机理。因此，目前对Si添加合金的研发，主要集中在提升其耐液态LBE腐蚀性能，很少从材料设计的角度兼顾其室温力学性能、高温性能、持久性能、蠕变性能和环境相容性等。同时，合金在液态LBE环境中的断裂韧性与疲劳性能等影响关键设备结构完整性与服役寿命的力学性能数据仍极度缺乏，主要受限于高温液态LBE环境中应变与裂纹长度等原位测量技术不成熟。基于此，本文主要介绍了中国科学院金属研究所部分科研团队在液态LBE环境相容性评价技术、商用T91与316钢的液态LBE环境相容性评价以及经微合金化的Si增强型(9%~12%)Cr铁素体/马氏体钢与316不锈钢的研究进展。
1 铁素体/马氏体钢T91与不锈钢316的LBE相容性
图1
图1
液态Pb-Bi共晶(LBE)腐蚀损伤测试装置实物图
Fig.1
Photos of corrosion damage test apparatuses in liquid lead-bismuth eutectic (LBE)
(a) immersion test apparatus (b) Pt/air sensor
(c) slow strain rate/creep test apparatus (d) fatigue test apparatus
基于液态LBE腐蚀浸泡试验装置与液态LBE溶解氧电极及控氧系统，开展了铁素体/马氏体钢T91 (Cr 8.77、Si 0.36、Mo 0.9、C 0.11、Ni 0.17、Mn 0.4、Fe余量，质量分数，%)与不锈钢316 (C 0.021、Si 0.43、Mn 1.52、P 0.016、S 0.003、N 0.056、Ni 13.0、Mo 2.39、Cr 17.82、Fe余量，质量分数，%)的液态LBE腐蚀性能研究。图2[4]为T91在550℃、不同DO浓度液态LBE环境中浸泡1000 h后的氧化膜截面形貌。在饱和氧(1.38 × 10-3%)条件下，T91表面氧化膜厚度约为48 μm，分为外层疏松多孔的Fe3O4、内层致密的Fe-Cr尖晶石氧化物及发生Cr选择性氧化的内氧化区；当DO浓度降低至1.26 × 10-6%时，氧化膜厚度降低至约7 μm，分为外层Fe-Cr尖晶石氧化物与发生Cr选择性氧化的内氧化区；当DO浓度降低至1.41 × 10-8%时，表面难以生成氧化膜，发生了基体溶解与Pb-Bi侵入；当DO浓度继续降低至1.12 × 10-9%时，溶解腐蚀与Pb-Bi侵入进一步加剧。图3为不锈钢316在550℃、不同DO浓度液态LBE环境中浸泡1000 h后的氧化膜截面形貌。在饱和氧条件下，316不锈钢表面氧化膜厚度约为14 μm，分为外层疏松多孔的Fe-Pb氧化物与Fe3O4以及内层致密的Fe-Cr尖晶石氧化物，氧化膜中存在大量的孔洞，有诱发裂纹导致氧化膜剥落的倾向；当DO浓度降低至1 × 10-7%时，表面几乎未观察到氧化膜，发生了基体溶解与Pb-Bi侵入。
图2
图2
铁素体/马氏体钢T91在550℃液态LBE中浸泡1000 h后的截面形貌[4]
Fig.2
Cross-section morphologies of oxide films on T91 steel after 1000 h exposure in liquid LBE at 550oC and different mass fractions of dissolved oxygen (DO)[4] (OOL—outer oxide layer, IOL—inner oxide layer, IOZ—inner oxide zone, OL—oxide layer)
(a) saturated oxygen (b) 1.26 × 10-6% DO (c) 1.41 × 10-8% DO (d) 1.12 × 10-9% DO
图3
图3
不锈钢316在550℃液态LBE中浸泡1000 h后的截面形貌
Fig.3
Cross-sectional morphologies of oxide films on 316 stainless steel after 1000 h exposure in liquid LBE at 550oC
(a) saturated oxygen (b) 1 × 10-7% DO
基于液态LBE腐蚀疲劳试验装置，开展了T91与316钢的液态LBE腐蚀疲劳性能研究。图4为T91与316钢在高温空气与饱和氧液态LBE环境中的低周疲劳寿命曲线。结果表明，T91在350℃空气中疲劳寿命位于ASME平均曲线附近；在350℃液态LBE环境中疲劳寿命急剧降低，且在高应变幅时下降得更为显著，在应变幅为1%时疲劳寿命已接近ASME设计曲线。316钢在400℃空气中疲劳寿命位于ASME平均曲线附近；在400℃液态LBE环境中，疲劳寿命与空气中的相当，仅在高应变幅(≥ 0.8%)条件下降低。图5为T91 (饱和氧，350℃，应变幅0.6%)与316钢(饱和氧，400℃，应变幅0.8%)在液态LBE环境中疲劳裂纹扩展区形貌。T91的疲劳断口表面未观察到疲劳辉纹，为典型的准解理开裂特征，发生了LME，导致疲劳寿命显著降低[16,17,19]；316钢的断口表面呈现疲劳辉纹特征，未发生LME，疲劳寿命与空气中的相当或略微降低。
图4
图4
T91和316钢在高温空气与液态LBE环境中疲劳寿命对比
Fig.4
Comparisons between fatigue life in air and liquid LBE (ASME—America Society of Mechanical Engineers)
(a) T91 steel (b) 316 stainless steel
图5
图5
T91和316钢在液态LBE环境中疲劳裂纹扩展区特征
Fig.5
Morphologies of fatigue crack propagation region in liquid LBE
(a) T91 steel, 350oC, strain amplitude 0.6%
(b) 316 stainless steel, 400oC, strain amplitude 0.8%
上述结果表明，控制液态LBE环境中的DO浓度，可提高T91钢的耐腐蚀性能，但316钢含有较高的Ni元素，在低氧条件下发生溶解腐蚀。兼顾T91与316钢耐腐蚀性能的最佳DO浓度仍需要更详细的实验数据来支撑选择。而且，铁素体/马氏体钢T91与不锈钢316在液态LBE环境中可能发生显著的LME或腐蚀降级，需优化其合金成分与制备工艺等来提高其固有的耐液态LBE腐蚀性能。
2 Si增强型铁素体/马氏体钢((9%~12%)Cr)和奥氏体不锈钢(316)
2.1 Si增强型铁素体/马氏体钢((9%~12%)Cr)
中国科学院金属研究所和中国原子能科学研究院在铁素体/马氏体钢T91 (ASME标准)成分的基础上，通过提高Si含量，并适当调整其他合金元素含量，研制出了Si增强型(9%~12%)Cr铁素体/马氏体钢，以下分别简记为9Cr-Si (9Cr系)和12Cr-Si (12Cr系)。2种Si增强型钢的耐LBE腐蚀性能得到明显提高，在550℃静态饱和氧的LBE环境下浸泡1000 h后9Cr-Si的氧化层厚度为24 μm (图6)，明显低于T91的48 μm (图2a)；12Cr-Si的氧化膜生长速率明显低于HT9 (图7a)，HT9和12Cr-Si在腐蚀10000 h后的氧化层厚度分别为75和42 μm (图7b和c)。此外，在550℃动态控氧(0.3 m/s，10-6%~10-7%)的LBE环境下，HT9和12Cr-Si的氧化层(1500 h)厚度分别为22和12 μm (图7d和e)。Si增强型(9%~12%)Cr铁素体/马氏体钢在静态和动态LBE环境下的优异耐腐蚀性能与其Si含量的提高有关：较高的Si含量在内氧化层与基体之间会形成富Si层(图6b和7e)，且Si倾向于在晶界处形成含Si氧化物颗粒，阻碍了元素沿晶界的快速扩散，从而提高了钢的耐LBE腐蚀性能[41]。
图6
图6
9Cr-Si在550℃静态饱和氧LBE环境下腐蚀1000 h后的氧化层形貌及元素分布
Fig.6
Cross-sectional back scattered electron (BSE) image (a) and corresponding line scanning (b) of 9Cr-Si alloy after exposure for 1000 h to stagn-ant oxygen-saturated LBE at 550oC
图7
图7
12Cr-Si系铁素体/马氏体钢LBE腐蚀实验结果
Fig.7
LBE corrosion test results of 12Cr-Si ferritic/martensitic steels
(a) corrosion layer thickness changes with time in stagnant oxygen-saturated LBE at 550oC (b, c) cross-sectional morphologies of HT9 steel (b) and 12Cr-Si steel (c) after exposure for 10000 h to stagnant oxygen-saturated LBE at 550oC (d, e) cross-sectional morphologies and corresponding EDS element mappings of HT9 steel (d) and 12Cr-Si steel (e) after exposure to flowing and oxygen controlled LBE for 1500 h (0.3 m/s, 10-6%-10-7%)
Si增强后可以显著提升(9%~12%)Cr铁素体/马氏体钢的耐LBE腐蚀性能，但Si作为一种强铁素体形成元素，也会提高高温δ铁素体的析出倾向[42]。如在铸态9Cr-Si中就存在约3.6% (面积分数)的δ铁素体(图8a)，该组织会降低(9%~12%)Cr铁素体/马氏体钢的冲击韧性、持久性能等。理论计算表明，9Cr-Si存在完全奥氏体化的温度区间，理论上可通过高温处理促使δ铁素体回溶，从而达到消除δ铁素体的目的。研究亦表明，铸态9Cr-Si经1150℃均质化处理后，可消除δ铁素体相(图8b)。均质化后的9Cr-Si经过热加工变形及正火和回火处理，得到回火马氏体组织(图9)，且在晶界和板条界析出了M23C6和NbC。因此，9Cr-Si获得了与T91钢相似的微观组织，其室温、高温拉伸性能以及韧脆转变温度(DBTT)与T91相当(图10)。值得注意的是，12Cr-Si在提高Si含量的同时，利用合金元素调控铬镍当量(CNB值)，结合制备工艺优化，使其δ-Fe面积分数< 1%，保证了合金具有较高的强度和韧性。
图8
图8
9Cr-Si铁素体/马氏体钢的铸态组织和均质化处理后的组织
Fig.8
SEM images of microstructures of as cast 9Cr-Si (a) and as-homogenized 9Cr-Si (b)
图9
图9
回火态9Cr-Si铁素体/马氏体钢的SEM像
Fig.9
SEM image of tempered 9Cr-Si ferritic/martensitic steel
图10
图10
9Cr-Si铁素体/马氏体钢和T91的室温、高温强度曲线及韧脆转变温度(DBTT)曲线
Fig.10
Strength curves at room temperature and high temperature (a) and ductile-to-brittle transition temperature (DBTT) curves (b) of 9Cr-Si and T91 steels
(9%~12%)Cr铁素体/马氏体钢在高温服役过程要求微观组织具有良好的稳定性和优异的持久性能，以保证核能系统结构的安全[43]。550℃时效3000 h实验结果表明，9Cr-Si的拉伸性能变化不大，其冲击韧性虽有一定程度的下降(图11a)，时效后合金钢的室温冲击功(114 J)大于100 J，仍具有较好的冲击韧性。继续延长时效时间至10000 h，钢的冲击功基本保持不变。研究表明，9Cr-Si的冲击功下降是由于析出了脆性的Laves相(图11b和c)，会引起局部应力集中。上述结果表明，Si增强后的9Cr-Si具有较好的组织稳定性。650℃下的高温持久实验测试结果(图12)表明，12Cr-Si的持久性能明显优于HT9，优异的高温持久性能得益于12Cr-Si钢中较多的强化相和低δ-Fe含量。
图11
图11
9Cr-Si铁素体/马氏体钢550℃时效后DBTT曲线及时效3000 h后的TEM像与元素分布图
Fig.11
DBTT curves of tempered and aged 9Cr-Si steel at 550oC (a), TEM image after 3000 h aging (b), and the elements mappings of the area denoted by the rectangle in Fig.11b (c)
图12
图12
650℃下12Cr-Si和HT9持久性能的对比
Fig.12
Creep-rupture strength of 12Cr-Si and HT9 steels at 650oC
综上所述，研发的Si增强型(9%~12%)Cr铁素体/马氏体钢，除了具有优异的耐LBE腐蚀性能，同时还有良好的室温、高温拉伸性能和冲击韧性，具有优异的持久性能和高温长时时效组织稳定性，有望满足铅冷快堆的设计要求。
2.2 Si增强316不锈钢
中国科学院金属研究所和中国原子能科学研究院在316钢成分的基础上，通过提高Si含量，并适当调整其他合金元素含量，成功研制出了Si增强奥氏体不锈钢，以下简记为ASS-Si。图13为ASS-Si和316钢经550℃饱和氧LBE腐蚀5000 h后氧化膜厚度随时间的变化曲线。可见，2种奥氏体钢的氧化膜厚度随时间的变化呈抛物线规律，其速率常数分别为0.18和0.44 μm2/h。ASS-Si在LBE中的腐蚀速率显著小于316钢。ASS-Si经饱和氧Pb-Bi腐蚀1000 h后氧化膜的截面形貌和元素分布如图14所示。表面形成了双层结构的氧化膜，外层为Fe3O4(图14b)，内层为Fe-Cr尖晶石，同时观察到Si的明显富集(图14c)。富Si氧化物具有较好的稳定性，提高了内层氧化膜的致密性，可抑制金属元素的向外扩散和O的内扩散，是ASS-Si耐LBE腐蚀性能提升的主要原因。
图13
图13
ASS-Si和316钢经550℃饱和氧LBE腐蚀5000 h后氧化层厚度随时间的变化曲线
Fig.13
Thicknesses of oxide scale of ASS-Si and 316 austenitic steels after exposure up to 5000 h to stagnant oxygen-saturated LBE at 550oC
图14
图14
ASS-Si经550℃饱和氧LBE腐蚀1000 h后氧化膜的截面形貌和元素分布
Fig.14
Cross-sectional BSE image (a) and corresponding EDS analyses (b,c) of ASS-Si austenitic steel after exposure for 1000 h to stagnant oxygen-saturated LBE at 550oC
进一步考察了流动控氧条件下ASS-Si和316钢的耐LBE腐蚀性能。图15为550℃、0.3 m/s流速、10-6%~10-7% DO浓度的液态LBE中腐蚀1500 h后的样品形貌。316钢表面出现了明显的溶解腐蚀，最大腐蚀深度约200 μm，破坏了样品的完整性(图15a)。样品表面未形成氧化膜，Pb-Bi已侵入基体，出现了Ni溶解导致的铁素体层(图15b)。相比之下，ASS-Si样品仍保持良好的完整性(图15c)。高倍照片显示，样品表面形成了约3 μm厚的连续、均匀的氧化膜(图15d)。可见，ASS-Si表面能形成保护性氧化膜，有效阻碍液态Pb-Bi的侵入和元素溶解，保证了其在流动控氧条件下的耐LBE腐蚀性能。
图15
图15
ASS-Si和316钢经550℃流动控氧LBE腐蚀1500 h后的截面形貌
Fig.15
Cross-sectional SEM images of 316 (a,b) and ASS-Si (c,d) austenitic steels after exposure for 1500 h to flowing oxygen-controlled LBE at 550oC (Figs.12b and d are enlarged views of the areas in Figs.11a and c, respectively)
然而，Si添加会影响奥氏体的稳定性，增加δ铁素体的形成倾向，进而损伤力学性能。与奥氏体相比，δ铁素体中Cr、Mo的过饱和度高、元素扩散速率快，高温服役过程中δ铁素体易分解为M23C6、Laves、σ等脆性相[44~46]。研究[47]表明，δ铁素体的分解速率比奥氏体中的相析出速率快2个数量级以上。另外，Si会促进奥氏体中σ相和G相((Ni,Fe,Cr)16(Nb,Ti)6Si7)的析出。中子辐照会进一步加速富Si相的析出[48]。为消除Si添加形成δ铁素体的不利影响，设计合金具有较高的C含量以调控Cr、Ni当量比。然而，550℃时效过程中，晶界处易析出M23C6。由于Si在碳化物中的溶解度较低，Si原子逐渐被排斥至周围基体中，增加了附近基体中Si的过饱和度。随着时效时间的延长，Si含量较高的G相和σ相从过饱和奥氏体基体中脱溶析出，导致奥氏体的稳定性降低，转变为α铁素体，最终在晶界附近形成了α铁素体+ G相+ σ相的混合组织(图16)。
图16
图16
含Si奥氏体钢经550℃时效处理1000 h后晶界附近的微观组织及元素分布
Fig.16
TEM image and the corresponding element mappings of Cr, Ni, Si, and Fe near grain boundaries of Si-modified austenitic steel after aging at 550oC for 1000 h
表1 ASS-Si奥氏体钢和Type 304奥氏体钢在550℃下的拉伸强度和持久强度 (MPa)
Table 1
| Steel | Yield strength | Tensile strength | 3000 h creep-rupture strength |
|---|---|---|---|
Type 304 ASS-Si | 106 201 | 349 433 | 210 220 |
3 总结与展望
(1) 铁素体/马氏体钢T91在550℃饱和氧液态LBE环境中发生快速氧化，DO浓度降低至1.26 × 10-6%时，腐蚀速率降低，DO浓度低于1 × 10-6%时，发生溶解腐蚀；T91的LME敏感性高，在350℃液态LBE环境中发生准解理开裂，低周疲劳寿命显著降低。
(2) 不锈钢316在550℃饱和氧液态LBE环境中发生快速氧化，DO浓度降低至1 × 10-7%时，发生溶解腐蚀，Ni选择性溶解，Pb-Bi侵入基体中；316不锈钢对LME不敏感，在400℃液态LBE环境中低周疲劳性能仅略微降低。
(3) Si增强型(9%~12%)Cr铁素体/马氏体钢表面生成富Si氧化膜，耐LBE腐蚀性能显著提高，具有良好的高温拉伸性能、组织稳定性和高温持久性能等。
(4) Si增强ASS-Si奥氏体钢中形成了稳定性更好的富Si氧化物，提高了氧化膜的致密性，在550℃静态饱和氧和动态控氧的液态LBE中具有良好的耐腐蚀性能；同时基于Nb / C质量比等关键元素控制，改善了Si添加对奥氏体稳定性带来的不利影响，使设计合金具有良好的力学性能。
(5) 对液态LBE环境中应用的Si增强型(9%~12%)Cr铁素体/马氏体钢及奥氏体不锈钢的优化，还需系统考核其可焊性与辐照性能，特别是辐照与液态LBE腐蚀耦合作用下材料的行为与机理。未来更应关注其高温长期服役过程中的组织、性能稳定性(如相析出机制)以及持久、蠕变与疲劳性能等，为铅冷快堆的安全运行和结构材料的迭代升级提供数据积累与技术支撑。
(6) LME是影响铅冷快堆结构材料综合性能的关键，现有模型难以解释LME发生的本质原因，未来应从原子尺度上揭示LME损伤机理，从而指导研发兼具耐LBE腐蚀与LME敏感性低的新型合金。
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图1
表1
