SiCf/SiC的蠕变性能与其纤维和基体等组元密切相关。三代SiC纤维具有良好的结晶度和化学计量比, 其复合材料的抗蠕变性能显著优于一代和二代SiCf/SiC复合材料[8], 如Hi-NicalonTM SiCf/SiC的1300 ℃/120 MPa蠕变断裂时间仅为0.83 h[11]。而采用三代SylramicTM-iBN纤维后, SiCf/SiC的抗蠕变性能显著提高, 1315 ℃/120 MPa蠕变断裂时间超过180 h[9]。对于不同基体的SiCf/SiC来说, 化学气相渗透(CVI)SiC基体具有优异的抗蠕变性能, 但CVI SiCf/SiC的孔隙率较高, 其抗蠕变性能比熔体渗透(MI)制备的SiCf/SiC(MI SiCf/SiC)差。而MI SiCf/SiC基体中的游离硅又限制了其最高使用温度。前驱体渗透裂解(PIP)制备的SiCf/SiC(PIP SiCf/SiC)消除了游离硅, 但是由于PIP SiC基体中存在严重微裂纹, PIP SiCf/SiC的抗蠕变性能远比CVI SiCf/SiC和MI SiCf/SiC差。SiCf/SiC的抗蠕变性能还与温度和应力等外部因素有关, 如CVI Hi-NicalonTM SiC/SiC的蠕变机制随应力和温度的不同而变化, 当蠕变应力远低于比例极限应力(σPLS)时, 其1315 ℃/69 MPa的蠕变断裂时间大于300 h[12⇓-14], 蠕变行为由CVI SiC基体控制; 当应力大于σPLS时, 材料在蠕变过程中发生基体开裂等损伤, 氧化性气氛进入材料内部造成氧化损伤, 材料的蠕变断裂时间显著缩短[12], 蠕变行为主要由纤维控制。
目前我国已经突破了碳化硅纤维和SiCf/SiC制备的关键技术。为了推进SiCf/SiC在我国航空发动机的应用, 本工作研究了国产二代Cansas-II SiC纤维平纹编织布增强的SiCf/SiC(2D-SiCf/SiC)在空气中的拉伸蠕变行为, 蠕变温度为1200~1400 ℃, 蠕变应力为80~140 MPa, 获得了应力和温度对蠕变性能的影响规律, 并利用扫描电子显微镜(SEM)观察蠕变断口形貌, 进一步揭示了2D-SiCf/SiC的蠕变损伤机理。
1 实验方法
1.1 实验材料
表1 国产二代Cansas-II碳化硅纤维的基本性能
Table 1
| Diameter/ μm | Density/ (g·cm-3) | Tensile strength/GPa | Tensile modulus/GPa |
|---|---|---|---|
| 14 | 2.74 | 2.7 | 270 |
表2 2D-SiCf/SiC的拉伸性能
Table 2
| Temperature/℃ | E/GPa | σPLS/MPa | σUTS/MPa | ε/% |
|---|---|---|---|---|
| RT | 273 | 115 | 282 | 0.57 |
| 1200 | 259 | 110 | 249 | 0.74 |
| 1300 | 223 | 92 | 229 | 0.60 |
σUTS: Ultimate Tensile Strength
1.2 实验方法
使用电子万能试验机(DDL 100, 中机试验装备股份有限公司)测试2D-SiCf/SiC的室温拉伸性能, 试样形状和尺寸如图1所示。使用应变片(BE120- 10AA, 中航电测仪器股份有限公司)测量应变, 加载速率为0.5 mm/min, 得到2D-SiCf/SiC的拉伸应力-应变曲线, 从而计算出弹性模量(E)、强度极限(σUTS)、σPLS和断裂应变(ε)等性能。蠕变试验在电子蠕变试验机(RDL50, 中机试验装备股份有限公司)上进行。高温拉伸蠕变试样如图1所示。蠕变试验在大气环境中进行, 蠕变温度为1200、1300和1400 ℃, 蠕变应力为80~140 MPa。对于未发生蠕变断裂试样, 开展室温拉伸试验, 测量其剩余拉伸性能。
图1
图1
拉伸和拉伸蠕变试样形状与尺寸(单位: mm)
Fig. 1
Dimensions and shape of tensile and tensile creep specimen (unit: mm)
利用SEM(TESCAN MIRA3, 捷克)观察试样的蠕变断裂断口形貌。切割试样近断口部分制成金相试样, 利用SEM观察纤维和基体损伤。利用SEM附带的电子能谱仪(EDS, OXFORD X-MaxN, 英国)分析纤维和基体的成分。
2 实验结果
2.1 蠕变曲线
图2显示了2D-SiCf/SiC的1200、1300和1400 ℃的典型蠕变曲线, 由图可见2D-SiCf/SiC 的蠕变曲线可分为减速蠕变和稳态蠕变两个阶段。随应力增加或温度升高, 稳态蠕变阶段缩短甚至会消失。
图2
图2
2D-SiCf/SiC的高温拉伸蠕变曲线
Fig. 2
Tensile creep curves of 2D-SiCf/SiC at high temperatures
(a) 1200 ℃, different stress; (b) 1300 ℃, different stress; (c) 1400 ℃, different stress; (d) Different temperatures, 100 MPa
在相同温度下, 随应力增加, 蠕变断裂时间显著缩短, 如图2(a~c)所示: 当蠕变温度为1200 ℃, 蠕变应力为110 MPa时, 2D-SiCf/SiC的蠕变断裂时间超过560 h; 当应力增大至140 MPa时, 蠕变断裂时间不到1 h。同时, 在相同应力下, 随温度升高, 蠕变断裂时间缩短, 蠕变断裂应变增加。当蠕变应力为100 MPa时, 2D-SiCf/SiC的1200 ℃蠕变断裂时间超过560 h, 蠕变断裂应变小于0.2%; 当温度升高到1300 ℃时, 蠕变断裂时间缩短至70 h; 当温度达到1400 ℃时, 蠕变断裂时间不到2 h, 蠕变断裂应变增至0.57%, 如图2(d)所示。从图2还可以看出, 要延长2D-SiCf/SiC的蠕变断裂时间, 应在提高温度的同时, 降低蠕变应力。当蠕变温度由1200 ℃升高到1300 ℃, 蠕变应力由110 MPa降低到90 MPa时, 2D-SiCf/SiC的蠕变断裂时间超过500 h。进一步升高温度到1400 ℃, 即使蠕变应力降低至80 MPa, 蠕变断裂时间也仅为80 h。从图2还可看出, 在 1200和1300 ℃下, 蠕变断裂应变随应力升高, 表现出先减小后增大的趋势。文献研究表明: 在相同的蠕变时间内, 应力增加会提高蠕变速率, 2D- SiCf/SiC的蠕变断裂应变也随之升高; 当应力过高, 蠕变断裂时间大幅缩短, 导致应变降低[11⇓⇓-14]。因此, 蠕变断裂应变的减小与增大受应力和蠕变断裂时间的影响。蠕变断裂应变先减小是由于蠕变断裂时间大幅缩短(从大于500 h降至几十小时), 后增大是因为应力增加导致蠕变速率增大了1~2个数量级。
根据图2所示的蠕变曲线, 可得到稳态蠕变速率、蠕变断裂应变, 蠕变断裂时间等性能参数, 见表3。可以看出, 稳态蠕变速率随温度升高或应力增加而增大, 且蠕变断裂时间越长, 稳态蠕变速率越小。稳态蠕变速率在1×10-10~5×10-10 /s之间时, 蠕变断裂时间超过500 h; 稳态蠕变速率约为10-9 /s左右时, 蠕变断裂时间为10~100 h; 稳态蠕变速率大于10-8 /s时, 蠕变断裂时间不到10 h。表3同时列出了日本朱世杰[13]和美国GE[15]等获得的MI-SiCf/SiC和CVI-SiCf/SiC两种国外二代材料在相近条件下的蠕变性能。可见, 国产2D-SiCf/SiC的稳态蠕变速率与两种国外材料接近, 但是其蠕变断裂时间显著少于MI-SiCf/SiC。
表3 2D-SiCf/SiC复合材料在空气中的高温拉伸蠕变性能
Table 3
| Specimen | Tempera- ture/℃ | Stress/ MPa | Steady-state creep strain rate/s-1 | Creep rupture time/h | Creep strain/% |
|---|---|---|---|---|---|
| 2D-SiCf/SiC | 1200 | 100 | 1.86×10-10 | >560 | 0.20 |
| 1200 | 110 | 4.95×10-10 | >560 | 0.23 | |
| 1200 | 120 | 2.16×10-9 | 42 | 0.18 | |
| 1200 | 140 | 8.45×10-7 | 0.1 | 0.21 | |
| 1300 | 90 | 5.62×10-10 | >500 | 0.28 | |
| 1300 | 100 | 1.73×10-9 | 70 | 0.20 | |
| 1300 | 120 | 1.86×10-8 | 6 | 0.28 | |
| 1400 | 80 | 1.32×10-9 | 88 | 0.33 | |
| 1400 | 100 | 1.46×10-8 | 1.6 | 0.57 | |
| CVI-SiCf/SiC[13] | 1300 | 75 | 2.3×10-9 | 111.1 | — |
| 1300 | 90 | 6.2×10-8 | 33.4 | — | |
| 1300 | 120 | 3.6×10-7 | 0.83 | — | |
| 1300 | 150 | 3.0×10-6 | 0.14 | — | |
| MI-SiCf/SiC[15] | 1204 | 125 | 3.0×10-11 | >1000 | 0.25 |
| 1204 | 140 | 2.3×10-10 | >1000 | 0.31 | |
| 1204 | 150 | — | 82.1 | 0.29 |
2.2 剩余拉伸性能
对于未发生蠕变断裂的2D-SiCf/SiC, 进一步测试其室温拉伸性能, 获得了图3所示的剩余拉伸应力-应变曲线, 其剩余拉伸性能见表4。可见, 不同条件蠕变后, 2D-SiCf/SiC的剩余拉伸性能显著低于原始材料。随着蠕变应力增加或温度提高, 2D-SiCf/SiC的剩余拉伸性能保持率降低。1200 ℃/100 MPa蠕变560 h后, 2D-SiCf/SiC的σUTS保持率为75%; 若蠕变应力增大到110 MPa, σUTS保持率降低至61%, σPLS、E和ε降低至原始材料的77%、77%和44%。当温度升高至1300 ℃, 蠕变应力降低至90 MPa时, 2D-SiCf/SiC的拉伸性能保持率约为50%, 显著低于2D-SiCf/SiC在1200 ℃蠕变后的拉伸性能保持率, 说明蠕变温度对该材料的蠕变损伤影响更大。
图3
图3
不同状态的2D-SiCf/SiC的拉伸应力-应变曲线
Fig. 3
Tensile stress-strain curves of as-received and crept 2D-SiCf/SiC
表4 2D-SiCf/SiC蠕变后的剩余拉伸性能
Table 4
| σUTS/MPa | E/GPa | σPLS/MPa | ε/% | ||
|---|---|---|---|---|---|
| As-received | 282 | 273 | 115 | 0.57 | |
| Crept | 1200 ℃/100 MPa | 211 | — | — | — |
| 1200 ℃/110 MPa | 173 | 209 | 89 | 0.25 | |
| 1300 ℃/90 MPa | 138 | 157 | 55 | 0.17 | |
2.3 断口形貌
为进一步揭示2D-SiCf/SiC的蠕变机理, 观察了典型试样的断口形貌, 根据断口上的特征区域判断蠕变损伤。图4(a)为2D-SiCf/SiC的室温拉伸断口形貌, 整个断口表面均能看到纤维拔出。图4(b)为1200 ℃/110 MPa蠕变560 h后2D-SiCf/SiC的室温拉伸断口形貌, 断口大部分区域的特点与图4(a)相似, 能从断口上看到明显的纤维拔出, 说明对应的区域基本没有发生氧化, 为瞬断区; 在该试样断口的边缘和穿刺纤维附近, 断口平齐, 基本无纤维拔出现象, 为氧化区[16]。从两类区域面积大小可以判断,蠕变过程中材料内部没有发生明显氧化, 仅在试样边缘和穿刺纤维附近发生氧化。图4(c, d)是蠕变温度为1200 ℃, 蠕变应力分别为120和140 MPa下2D-SiCf/SiC蠕变断裂后的断口形貌。蠕变应力为120 MPa时, 蠕变断裂时间为42 h, 其断口的氧化区面积显著增大。瞬断区面积缩小, 该区域上的纤维拔出长度明显增加(图4(c))。蠕变应力为140 MPa时, 蠕变断裂时间仅为0.1 h, 这时断口上也能够看到氧化区和瞬断区, 其氧化区面积小于图4(c)。对比图4(b~d)可知, 不同蠕变时间导致的蠕变断口呈现不同的断口特点, 即蠕变时间越长, 氧化区面积越大。当蠕变应力较小时, 材料的基体开裂程度低, 到达材料内部的氧含量少, 即使蠕变时间大于560 h, 材料内部也没有发生明显氧化。蠕变应力增加, 基体开裂程度增大, 氧气能够通过基体裂纹或穿刺纤维等进入材料内部, 使得材料内部氧化程度增加。蠕变温度升高, 其断口宏观形貌也呈现类似特点。
图4
图4
2D-SiCf/SiC在不同条件下的断口宏观形貌
Fig. 4
Macroscopic morphologies of fractures of 2D-SiCf/SiC under different conditions
(a) As-received; (b) Crept at 1200 ℃/110 MPa for 560 h; (c) Crept at 1200 ℃/120 MPa; (d) Crept at 1200 ℃/140 MPa
前面结果可知2D-SiCf/SiC蠕变断裂断口氧化区的形成与蠕变过程密切相关, 蠕变过程中材料的氧化与蠕变温度、应力和时间等密切相关, 因此可从氧化区大小和形貌判断蠕变损伤过程。图5显示了2D-SiCf/SiC的3个典型蠕变断口的氧化区形貌, 其蠕变断裂时间分别是42、70和88 h。可以看出: 蠕变断口氧化区上没有纤维拔出现象, 纤维拔出长度短意味着纤维/基体界面结合强度高。2D-SiCf/SiC的不同温度断口氧化区形貌不同。1200 ℃/120 MPa蠕变断口氧化区表面形成氧化物(图5(a)), 据EDS分析可知, 这些氧化物主要为SiO2。蠕变温度升高至1300 ℃, 氧化程度加剧, SiO2熔体覆盖在材料断口氧化区表面, 同时可见图5(b)所示的气体溢出气泡。蠕变温度升高到1400 ℃时, 材料的氧化程度进一步加剧, 形成更多SiO2熔体覆盖于断口氧化区表面, 如图5(c)所示。因此蠕变温度越高或蠕变断裂时间越长, 蠕变断口表面的氧化程度越大。
图5
图5
2D-SiCf/SiC在不同条件蠕变断裂的断口氧化区形貌
Fig. 5
Morphologies of fracture oxidation zone of 2D-SiCf/SiC under different creep rupture conditions
(a) 1200 ℃/120 MPa; (b) 1300 ℃/100 MPa; (c) 1400 ℃/80 MPa
蠕变过程中, 界面处的BN氧化成B2O3气体逐渐挥发会造成纤维/基体界面间隙[17]。部分B2O3熔体可与基体或纤维氧化生成的SiO2融合形成了硼硅酸盐熔体。随着基体或纤维氧化程度增大, 硼硅酸盐熔体中SiO2含量增加, 其黏度增大, 最终固化为高SiO2含量的玻璃[18], 从而造成界面粘结, 如图6(a)所示。界面粘结将导致纤维/基体界面结合强度变大, 裂纹无法偏转, SiO2的填充也会增大应力集中程度, 进而使蠕变裂纹扩展, 造成蠕变断裂。此外, 在1200和1300 ℃蠕变断裂后, 还发现图6(b)所示的纤维氧化。这是因为在这两个温度下SiC的氧化速率低于1400 ℃的氧化速率。研究表明: 1300 ℃条件下, SiC氧化生成SiO2的速率为3.76×10-2μm2/h [19], 低于1400 ℃的氧化速率(5.89×10-2μm2/h)。没有生成足够的SiO2填充全部纤维/基体界面, 当BN界面被消耗形成空隙后, 纤维被氧化使其有效直径变小, 纤维强度降低, 进而造成纤维断裂。因此不论是界面粘结还是纤维氧化, 均会促使基体裂纹张开和扩展, 进而造成蠕变断裂。
图6
图6
1300 ℃/100 MPa 蠕变过程中2D-SiCf/SiC中发生的界面和纤维损伤
Fig. 6
Interface and fiber damage in 2D-SiCf/SiC during creeping at 1300 ℃/100 MPa
(a) SiO2 filled; (b) SiO2 unfilled
2.4 截面形貌
由前述结果可知, 2D-SiCf/SiC在1200 ℃/110 MPa和1300 ℃/90 MPa两种条件下均未发生蠕变断裂, 但是1300 ℃/90 MPa蠕变后, 复合材料的剩余拉伸性能显著下降。进一步分析两个试样的截面微观组织, 如图7所示, 1200 ℃/110 MPa蠕变560 h后, 纤维和界面保持相对完好, 复合材料内部没有看到明显的氧化; 1300 ℃/90 MPa蠕变500 h后, SiC纤维发生了开裂, 说明纤维在蠕变过程中发生了严重损伤, 其性能退化。这就是2D-SiCf/SiC 经1300 ℃/90 MPa蠕变500 h后的拉伸性能显著下降的原因。
图7
图7
2D-SiCf/SiC蠕变后的纵向纤维微观组织
Fig. 7
Longitudinal fiber microstructures of 2D-SiCf/SiC after crept
(a)1200 ℃/110 MPa for 560 h; (b)1300 ℃/90 MPa for 500 h
3 分析与讨论
实验结果显示2D-SiCf/SiC的蠕变行为和损伤模式与应力和温度密切相关, 当蠕变应力低于σPLS时, 2D-SiCf/SiC有良好的蠕变性能, 基体不会发生明显开裂, 氧气无法大量进入材料内部, 复合材料蠕变560 h后的拉伸断口也反映出这个特征。这时蠕变载荷由纤维和基体共同承担, 其蠕变行为由基体和纤维共同控制。当蠕变应力大于σPLS时, 基体发生明显开裂, 氧气通过基体裂纹进入材料内部, 使基体、纤维和界面发生氧化, 发生界面粘结和纤维氧化等损伤(如图6所示)。增加应力或提高温度, 基体开裂程度增大, 此时蠕变行为主要由纤维决定。更高的温度和应力造成纤维损伤程度加剧, 这与文献[9]的研究结果一致。
从表2可知, 随着温度升高, 2D-SiCf/SiC的σPLS下降。在相同应力下, 蠕变温度越高, σPLS下降越小, 2D-SiCf/SiC的基体开裂程度越大, 因此抗蠕变性能越差。2D-SiCf/SiC的σPLS与其使用的纤维和基体工艺密切相关[9]。本研究中2D-SiCf/SiC的σPLS与CVI Hi-Nicalon SiCf/SiC的接近[13], 但低于MI-SiCf/SiC的[15]。为进一步说明基体类型对于蠕变性能的影响, 这里使用Larson-Miller参数(LMP)估计和比较了不同基体类型SiCf/SiC的蠕变断裂时间[9]。LMP是将蠕变断裂时间与蠕变应力和温度建立联系的参数, 关系如公式(1)所示。对于SiCf/SiC, D为22[20]。图8显示了2D-SiCf/SiC在不同应力下的LMP参数, 为了对比还列出了部分不同基体和纤维的SiCf/SiC的数据[9,13,15]。从图中可以看出, 在相同条件下, 2D-SiCf/SiC的蠕变性能与其使用的纤维和基体类型密切相关。纤维代次越高, 2D-SiCf/SiC的抗蠕变性能越好。MI SiCf/SiC的抗蠕变性能优于CVI SiCf/SiC。
其中T是温度(K), tR是蠕变断裂时间, D是常数。
图8
4 结论
1) 在1200~1400 ℃温度范围内, 2D-SiCf/SiC的蠕变断裂时间随温度升高或应力增加而缩短。在1200 ℃/110 MPa条件下, 其蠕变断裂时间大于560 h, 最小蠕变速率为4.95×10-10 /s, 蠕变性能达到国外类似材料水平。
2) 2D-SiCf/SiC的蠕变断裂时间与氧化损伤程度密切相关。在1200 ℃/110 MPa条件下, 氧化区面积小, 蠕变断裂时间大于560 h。增加应力或提高温度会使氧化区面积增大, 蠕变断裂时间缩短。
3) 2D-SiCf/SiC在1200~1400 ℃的蠕变行为主要是基体开裂和氧化损伤。基体开裂后, BN界面氧化形成间隙, 被氧化的纤维直径明显减小, SiC氧化生成的SiO2填充界面导致界面粘结, 这进一步促使基体裂纹张开和扩展。
参考文献
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Polymer-derived method is the main preparation method for continuous SiC fiber. So far, three generations of SiC-based fibers have been developed. SiC fiber is mainly used as reinforcement for SiC matrix to prepare SiCf/SiC composite, which exhibits excellent properties such as high-temperature resistance, oxidation resistance, high-temperature creep resistance, and neutron radiation stability. Therefore, it is considered as a promising preferred alternative material for structural components which works under harsh environments. In this review, the domestic and oversea development of the polymer-derived SiC fibers are introduced, in particular the different approaches for preparation and the characteristics of the 3rd generation fibers. Then, the development of SiCf/SiC preparation technology and properties are reviewed and the relationship between the innovation in fabrication of SiCf/SiC composite and the development of SiC fiber is hightlighted. Finally, the application progress of SiCf/SiC composite in high performance aero-engines and fusion power cores is summerized. Meanwhile, some perspectives on the development of domestic SiC fibers and SiCf/SiC are also discussed.
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