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金属学报, 2026, 62(3): 445-457 DOI: 10.11900/0412.1961.2024.00106

研究论文

GH4151难变形高温合金的拉伸行为及其断裂失效机制

崔天亮1,2, 谢兴飞,1,2,3, 温晓灿1,2,3, 吕少敏1,2,3, 曲敬龙,1,2,3, 杜金辉1,2,3

1.钢铁研究总院 高温材料研究所 北京 100081

2.北京钢研高纳科技股份有限公司 北京 100081

3.四川钢研高纳锻造有限责任公司 德阳 618000

Tensile Behavior and Fracture Mechanism of Hard-to-Deform GH4151 Superalloy

CUI Tianliang1,2, XIE Xingfei,1,2,3, WEN Xiaocan1,2,3, LYU Shaomin1,2,3, QU Jinglong,1,2,3, DU Jinhui1,2,3

1.High-Temperature Materials Institute, Central Iron and Steel Research Institute, Beijing 100081, China

2.Beijing Gaona Materials & Technology Co. Ltd., Beijing 100081, China

3.Sichuan Gaona Forging Co. Ltd., Deyang 618000, China

通讯作者: 曲敬龙，qujinglong@cisri.cn，主要从事先进航空发动机涡轮盘材料制备技术研究；谢兴飞，xiexingfei@cisri.com.cn，主要从事变形高温合金研究

责任编辑: 梁烨

收稿日期: 2024-04-08   修回日期: 2024-11-25  

基金资助: 国家科技重大专项项目(J2019-VI-0006-0120)
国家自然科学基金项目(52274330)
国家自然科学基金项目(52074092)

Corresponding authors: QU Jinglong, senior engineer, Tel: 13810256459, E-mail:qujinglong@cisri.cn;XIE Xingfei, senior engineer, Tel: 18801928583, E-mail:xiexingfei@cisri.com.cn

Received: 2024-04-08   Revised: 2024-11-25  

Fund supported: National Science and Technology Major Project(J2019-VI-0006-0120)
National Natural Science Foundation of China(52274330)
National Natural Science Foundation of China(52074092)

作者简介 About authors

崔天亮，男，1996年生，博士生

摘要

为探明服役过程中温度和应力快速叠加造成的涡轮盘损伤，本工作以涡轮盘用难变形GH4151合金为例，采用SEM、TEM、EDS和EPMA等表征手段，研究了不同温度下GH4151合金的拉伸行为和显微组织演化规律，分析了拉伸断裂失效机制。结果表明，随实验温度升高，GH4151合金的塑性呈现出先减小后增大的趋势，断裂方式由混合断裂逐渐向沿晶断裂转变。共格γ/γ′界面和MC界面位错塞积是导致混合断裂产生的主要原因，晶界位错塞积加速了高温高应力下O原子向晶界弹性应力场或缺陷处的富集，导致晶界动态脆化，引起沿晶断裂，造成GH4151合金在650~800 ℃高温拉伸时塑性下降。950 ℃拉伸时，GH4151合金强度迅速下降，裂纹以较慢的速率扩展，导致塑性增加。

关键词： 镍基高温合金; 拉伸性能; 断裂机制; 变形机制; 显微组织

Abstract

GH4151 is a heavy alloy, hard-to-deform Ni-based superalloy with service temperatures reaching up to 750-800 oC. It is an important candidate material for high-temperature alloys used in the turbine disks of the new-generation advanced aeroengines. During service, the rapid superposition of temperature and stress makes turbine disks susceptible to damage. This study explores the use of hard-to-deform GH4151 alloy used for turbine disks. The tensile behavior of the GH4151 alloy within a temperature range of 23-950 oC was investigated using advanced techniques such as SEM, TEM, EDS, and EPMA. The microstructural changes, deformation microstructure, and their impact on the fracture mechanism were analyzed, and the fracture failure mechanisms of the alloy at various temperatures were elucidated. The results indicate that yield strength and tensile strength initially decrease gradually, followed by a rapid decline with increase in experimental temperature. Meanwhile, elongation after fracture of the alloy decreased initially and increased with increasing experimental temperature. The fracture mode transitioned from a mixed fracture to an intergranular fracture. Further research showed that during tensile testing at temperatures of 23-550 oC, deformation primarily occurred in the γ channels, with a significant accumulation of dislocations at the γ/γ′ interfaces. This led to the tearing of the γ/γ′ interfaces and the formation of microvoids, which in turn generated a transgranular fracture. The intergranular fracture within the mixed-fracture mode is attributed to the stress concentration at MC carbide interfaces, resulting in the formation of voids. During tensile testing at temperatures of 650-800 oC, cracks were initiated via an intergranular fracture, and propagated through mixed-fracture modes. Deformation occurred simultaneously in the γ channels and the γ′ phase. Dislocation pile-up at the grain boundaries accelerated the enrichment of O atoms toward the elastic stress fields or the enrichment of defects at the grain boundaries under high-temperature and high-stress conditions. Such enrichment led to dynamic embrittlement of the grain boundaries, causing intergranular fracture, which reduced the elongation after fracture. As the strain increased, crack propagation was accelerated, reducing the time available for the O atoms to dynamically embrittle the grain boundaries. When the accumulation of dislocations at the γ/γ′ interfaces reached a critical value, crack propagation shifted to a mixed-fracture mode dominated by transgranular fracture. During tensile testing at 950 oC, cracks were initiated and propagated via intergranular fracture. The morphology of the γ′ phase changed to an approximately spherical shape, reducing the hindrance to dislocation motion. This reduction did not lead to the coalescence of microvoids at the γ/γ′ interfaces, and thus, no transgranular fracture occurred in the samples tested at 950 oC. Because of the decrease in tensile strength at 950 oC, the external stress applied was reduced, and crack propagation slowed down, elongation was increased.

Keywords： Ni-based superalloy; tensile property; fracture mechanism; deformation mechanism; microstructure

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本文引用格式

崔天亮, 谢兴飞, 温晓灿, 吕少敏, 曲敬龙, 杜金辉. GH4151难变形高温合金的拉伸行为及其断裂失效机制[J]. 金属学报, 2026, 62(3): 445-457 DOI:10.11900/0412.1961.2024.00106

CUI Tianliang, XIE Xingfei, WEN Xiaocan, LYU Shaomin, QU Jinglong, DU Jinhui. Tensile Behavior and Fracture Mechanism of Hard-to-Deform GH4151 Superalloy[J]. Acta Metallurgica Sinica, 2026, 62(3): 445-457 DOI:10.11900/0412.1961.2024.00106

随着航空发动机的推重比不断提升，涡轮盘前进口温度不断提高，对涡轮盘用高温合金的承温能力也提出了更高的要求[1~3]。GH4151合金是一种高合金化难变形镍基高温合金，其沉淀强化元素Al + Ti + Nb含量超过10% (质量分数)，γ'相体积分数约为55%，通过优化合金成分以及调控γ'相分布，其力学性能和高温组织稳定性得到显著提升[4~7]，服役温度可达800 ℃，是新一代先进航空发动机涡轮盘用高温合金的重要材料[8,9]。涡轮盘工作条件非常苛刻，在发动机启停或快速模式切换过程中，温度和应力快速叠加，极易造成涡轮盘损伤[10]。因此，研究不同温度下GH4151合金的拉伸行为及其断裂机制，对GH4151合金工程化应用具有理论指导作用。

近年来，国内外学者对涡轮盘用变形高温合金的拉伸失效断裂行为做了相关研究。Duan等[11]研究表明，变形高温合金在室温拉伸时呈现混合断裂特征，750 ℃高温拉伸断口呈现沿晶断裂特征。Wan等[12]研究发现，合金的韧性断裂是由于应力作用下微孔洞合并所致，沿晶断裂是由晶界第二相引起的应力集中所致。Németh等[13]发现，随着温度升高，合金延伸率呈现出先下降后增加的趋势，这主要是环境辅助晶界氧化所致，同时晶界滑动缓解了晶界处的位错塞积，进而提高了高温阶段的延伸率。高温合金的塑性对实验温度具有较高的依赖性，晶界氧化[14]、晶界析出相[15]、滑移面脱黏[16]和晶界滑动[13]均会导致合金高温塑性下降，再结晶和晶界滑动[16]是引起高温塑性提升的主要原因。此外，Li等[5]研究表明，铸态GH4151高温合金裂纹的开裂机理为冷却过程中冷裂纹的形成。Jia等[17]对铸态GH4151合金的开坯裂纹进行了分析，发现析出相边缘及晶界附近细小孔洞是导致合金产生沿晶断裂的主要原因。李育升等[18]研究了热处理工艺对GH4151合金冲击性能的影响，认为晶界上连续或不连续分布的第二相导致合金发生脆性断裂。Pan等[19]认为，GH4151合金蠕变过程中沿晶断裂的产生与晶界MC碳化物和γ'相有关。Gai等[20]研究了固溶温度对GH4151合金拉伸性能的影响，发现随着固溶温度升高，晶粒尺寸增大，晶界强度降低，导致750 ℃拉伸断口由韧性断裂转变为混合断裂。作为航空发动机涡轮盘重要的备选材料之一，深入探究GH4151合金在不同温度阶段的断裂机理对其实际应用具有重要意义。然而，目前仍缺乏针对GH4151合金在不同温度下的拉伸行为及其断裂机制的系统性研究。

基于此，本工作研究了GH4151合金在23~950 ℃范围内的拉伸行为和显微组织演化，分析了不同温度下GH4151合金的拉伸断裂机制，以期更好地理解不同工作条件下涡轮盘材料的损伤和失效过程，从而为提高涡轮盘的可靠性和安全性提供理论依据。

1 实验方法

实验所用GH4151合金盘锻件的化学成分(质量分数，%)为：C 0.04，Co 15.0，Cr 11.0，Mo 4.5，W 3.0，Al 3.0，Ti 2.75，Nb 3.55，Ni余量。实验所用拉伸样品均沿弦向进行取样(图1a)，拉伸样品形状及尺寸如图1b所示。拉伸样品的平均晶粒度为ASTM 8级(依据ASTM E112标准)，晶粒尺寸为22.3 μm，如图1d所示。

图1

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图1   GH4151合金拉伸样品取样位置、尺寸和表征位置示意图以及初始组织OM像

Fig.1   Schematics of sampling position (a), tensile sample size (unit: mm) (b), and material characterization positions (c) and OM image of the initial microstructure (d) of GH4151 alloy


依据GB/T 4338—2006标准，采用WE300B拉伸试验机在23、200、400、550、650、700、750、800和950 ℃条件下对GH4151合金进行拉伸实验，同时在真空条件下进行800 ℃拉伸实验作为对比实验，所有拉伸实验的应变速率均保持一致(屈服前0.003 min-1、屈服后0.03 min-1)，每个条件下设置3个平行样品。利用电解抛光和电解腐蚀方法制备扫描电子显微镜(SEM)样品，电解抛光在20%H2SO4 + 80%CH3OH溶液(体积分数，下同)中进行，电压20 V，抛光时间20 s；电解腐蚀在30 mL H3PO4 + 3 g CrO3 + 2 mL H2SO4溶液中进行，电压5 V，腐蚀时间2 s。采用Tenupol-5型双喷电解设备制备透射电子显微镜(TEM)样品，双喷液为10%HClO4 + 90%C2H5OH溶液，温度-25 ℃，电流80 mA。为了防止引入O原子，采用机械抛光+超声波清洗的方法制备电子探针显微分析(EPMA)样品，实验前样品需在高温下(200 ℃)烘烤10 min。实验样品观测位置如图1c所示，采用GX71型光学显微镜(OM)分析断口表面裂纹萌生区和裂纹扩展区特征。采用配备能谱仪(EDS)和电子背散射衍射(EBSD)系统的JSM-7800 F场发射SEM分析γ'相和断口微观组织以及晶粒组织取向特征。采用JXA-8530F型EPMA定量分析合金元素的分布信息。采用G2 F20型TEM分析合金变形微观组织。采用JmatPro软件计算不同温度下GH4151合金的晶格错配度。采用Image J软件统计不同温度下GH4151合金中二次γ'相的球形度，每个温度分别在10个不同视场下统计200个二次γ'相的球形度。

2 实验结果

2.1 拉伸性能

图2为不同温度下GH4151合金的工程应力-应变曲线和拉伸性能。可以看出，当实验温度从23 ℃升高至650 ℃时，GH4151合金的屈服强度从1204 MPa下降至1113 MPa，抗拉强度从1628 MPa下降至1487 MPa。在700~800 ℃范围内，随着实验温度升高，屈服强度从1083 MPa下降至902 MPa，抗拉强度从1347 MPa下降至1080 MPa。当温度到达950 ℃时，屈服强度下降至332 MPa，抗拉强度下降至497 MPa。此外，当实验温度从23 ℃升高至800 ℃时，GH4151合金的断后伸长率从15%下降至7%，当温度继续升高至950 ℃时，合金断后伸长率提升至14%。

图2

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图2   不同温度拉伸后GH4151合金的工程应力-应变曲线和拉伸性能

Fig.2   Engineering stress-strain curves (a) and tensile properties (b) of GH4151 alloy after tensile testing at various temperatures


2.2 断口形貌

图3为23和550 ℃拉伸后GH4151合金拉伸断口横截面的OM和SEM像。如图3a~c所示，23 ℃拉伸样品中存在剪切唇、大量解理面以及少量浅韧窝和沿晶裂纹；其断裂模式为以穿晶断裂为主的混合断裂，裂纹在样品表面萌生并以混合断裂的方式扩展。如图3d~f所示，550 ℃拉伸样品中同样存在大量解理面以及少量浅韧窝和沿晶裂纹，而剪切唇完全消失；其断裂模式仍为以穿晶断裂为主的混合断裂，但塑性出现一定程度的下降，这与图2b所示的塑性的变化趋势一致。

图3

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图3   23和550 ℃拉伸后GH4151合金断口横截面的OM和SEM像

Fig.3   Cross-sectional OM (a, d) and SEM (b, c, e, f) images of fracture morphologies of GH4151 alloy after tensile testing at 23 oC (a-c) and 550 oC (d-f)


图4为650~800 ℃拉伸后GH4151合金断口横截面的OM和SEM像。可以看出，650~800 ℃拉伸样品断口呈现出的断裂特性一致，即：存在明显的裂纹萌生区域和裂纹扩展区域(如图4a~d中黑色虚线标识区域)。差异在于，随着拉伸温度的升高，裂纹萌生区域面积逐渐增大，裂纹扩展区域相应减小。为进一步分析650~800 ℃拉伸样品的断裂方式，选取了具有代表性特征的800 ℃拉伸样品进行研究。图4e为800 ℃拉伸样品裂纹萌生区域的组织形貌。该区域表现出明显的沿晶断裂特征，说明650~800 ℃条件下样品裂纹萌生区域的断裂方式为沿晶断裂。g和区域表现出明显的沿晶为800 ℃拉伸样品裂纹扩展区域组织形貌。断口具有明显的解理面、穿晶撕裂棱和少量的沿晶裂纹，此区域的断裂模式为以穿晶断裂为主的混合断裂模式。因此，在650~800 ℃拉伸样品中，裂纹以沿晶断裂方式萌生，并以混合断裂模式扩展。值得注意的是，裂纹萌生区和裂纹扩展区具有明显的颜色差别，这可能与高应力氧化气氛下的环境损伤有关[13]。此外，550和800 ℃拉伸样品的混合断裂区域存在少量破碎或断裂的小平面(图3f中点1和图4f中点2位置)。采用EDS分析断裂小平面的化学成分，结果如表1所示。可以看出，这些小平面为富Nb、Ti的MC碳化物，这些MC碳化物可能与混合断裂中的沿晶断裂有关。

图4

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图4   650~800 ℃拉伸后GH4151合金断口横截面的OM和SEM像

Fig.4   Cross-sectional OM (a-d) and SEM (e-g) images of fracture morphologies of GH4151 alloy after tensile testing at 650 oC (a), 700 oC (b), 750 oC (c), and 800 oC (d-g) (The regions enclosed by the black dotted lines in Figs.4a-d represent the crack initiation areas, with the exterior demarcated as the crack propagation areas) (e) crack initiation zone (f, g) crack propagation zones


表1   图3f和4f中点1和2化学成分的EDS分析结果 (mass fraction / %)

Table 1  EDS analysis results of chemical compositions of points 1 and 2 in Figs.3f and 4f

PointCAlTiCrCoNiNbMoW
18.790.4511.714.675.7819.1941.144.264.01
28.051.089.136.359.8620.2736.833.584.85

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当温度达到950 ℃时，样品从边缘到心部已经完全被氧化，断口表面形成了一层氧化膜(图5a)。断口处存在明显的沿晶裂纹，且不存在解理面和韧窝(图5b和c)。因此，在950 ℃拉伸样品中，裂纹主要以沿晶断裂方式萌生并扩展。

图5

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图5   950 ℃拉伸后GH4151合金断口横截面的OM和SEM像

Fig.5   Cross-sectional OM (a) and low (b) and high (c) magnified SEM images of fracture morphologies of GH4151 alloy after tensile testing at 950 oC


2.3 微观组织

图6为不同温度拉伸后GH4151合金断口处二次γ'相的SEM像。可见，当拉伸温度为23~800 ℃时，二次γ'相形貌变化较小；当拉伸温度达到950 ℃时，部分二次γ'相形貌转变为近似球形。球形度(Φ)是一种基于二维图像分析的形状因子定义方法，其计算公式为[21]：

Φ=4πAP2

式中，A为颗粒投影面积，P为颗粒投影周长。Φ越接近于1，颗粒越接近于球形。如表2所示，拉伸温度低于800 ℃时，二次γ'相的平均球形度仅发生轻微变化；拉伸温度为950 ℃时，二次γ'相的平均球形度显著提高，由800 ℃时的0.808提高至0.849。实验结果(图6)表明，950 ℃拉伸样品中二次γ'相已发生球形化。二次γ'相的形貌与晶格错配度有关，晶格错配度的绝对值越接近于0.4，γ'相的形貌越接近立方形；晶格错配度的绝对值越接近于0，γ'相形貌接近于球形[22]。通过JmatPro软件计算得出，800和950 ℃拉伸后，GH4151合金晶格错配度的绝对值分别为0.29和0.01。综上所述，拉伸温度低于800 ℃时，GH4151合金具有较强的组织稳定性[19,23]，从而保持了二次γ'相尺寸和形貌的稳定。当温度升高至950 ℃后，晶格错配度的绝对值下降，γ'相形貌向球形转变。

图6

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图6   不同温度拉伸后GH4151合金断口处二次γ'相的SEM像

Fig.6   SEM images of secondary γ' phases at fractures of GH4151 alloys after tensile testing at 23 oC (a), 550 oC (b), 650 oC (c), 750 oC (d), 800 oC (e), and 950 oC (f)


表2   不同温度拉伸后GH4151合金断口中二次γ'相球形度

Table 2  Circularities of secondary γ' phases in GH4151 alloys after tensile testing at various temperatures

Tensile temperatureoCMedian circularityAverage circularity
230.7930.782
5500.7870.785
6500.8060.796
7500.7890.786
8000.8110.808
9500.8520.849

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图7为23、750和950 ℃拉伸后GH4151合金的TEM分析。如图7a~c所示，23 ℃拉伸后，γ通道内存在高密度位错和位错对切割γ'相形成的少量反相畴界(antiphase boundary，APB)。由于共格应变强化作用，1/2<110>位错被γ/γ'界面产生的共格应变场吸附，位错被限制在γ通道内，仅有少量γ'相被切割。因此，23 ℃拉伸时塑性变形主要发生在γ通道内。

图7

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图7   不同温度拉伸后GH4151合金变形组织的TEM分析

Fig.7   Bright field (BF) (a, c, d, g-i) and dark field (DF) (b, e) TEM images, and SAED pattern (f) of GH4151 alloy after tensile fracture at various temperatures (MT—microtwinning) (a-c) 23 oC (d-f) 750 oC (g-i) 950 oC


如图7d和e所示，750 ℃拉伸样品中存在穿过基体和γ'相的扩展层错(extended stacking fault，ESF)，以及单独存在于γ'相内部的孤立层错(isolated stacking fault，ISF)。对图7e中带状组织进行选区电子衍射(SAED)分析，结果如图7f所示。该带状组织为微孪晶(microtwinning，MT)。结果表明，750 ℃拉伸实验过程中，基体和γ'相同时发生变形。随着温度升高，共格应变强化效果减弱，位错可从共格应变场中逃逸并发生分解，如 式(2)所示[24]：

a2<101> =a3<211>+a6<121>

式中，a为晶格常数。a / 6<112>不全位错剪切γ基体和γ'相形成扩展层错；a / 6<112>不全位错连续剪切{111}面，经过原子重排消除高能Al—Al键后形成微孪晶；a / 3<112>不全位错切割γ′相产生孤立的层错[24]。如图7g~i所示，950 ℃拉伸样品中位错和微孪晶数量显著增加。在此阶段，合金晶格错配度降低，γ′相形貌转变为近球形，γ/γ′共格界面的应变强化效果减弱，其对位错阻碍作用也随之降低。高温下合金层错能增加，不全位错易束集成全位错并发生交滑移或攀移，因此合金中位错密度增大。随着实验温度升高，热激活效应导致原子扩散速率增加，促进了层错原子重新排列，从而使得微孪晶密度增加。

3 分析与讨论

3.1 23~550 ℃拉伸过程中合金的断裂机制

在23~550 ℃温度范围内，GH4151合金的断裂模式为以穿晶断裂为主的混合断裂。由韧性撕裂棱和解理面主导的穿晶断裂可能与二次γ′相脱黏有关[12,25,26]。如图8所示，解理面和韧性撕裂棱前沿存在纳米级浅韧窝，这些浅韧窝尺寸与二次γ′相接近，因此浅韧窝可能是二次γ'相脱黏遗留的产物。图9为23~550 ℃拉伸过程中GH4151合金穿晶断裂示意图。γ/γ′相间存在着微小的间隙，间隙数量随着析出相数量增加而增加[27]。值得注意的是，GH4151合金具有较高的γ′相含量，体积分数约为55%。此阶段大部分位错被限制在γ通道，塑性变形主要发生在γ基体中，而二次γ′相基本不发生变形。随着应变的累积，位错在γ/γ′界面大量积聚，当γ基体应变累积到临界值时，γ基体与二次γ′相间塑性变形不协调，导致细小的间隙撕裂形成微孔洞，引起二次γ'相脱黏，进而导致穿晶断裂。

图8

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图8   550 ℃拉伸后GH4151合金断口表面的SEM像

Fig.8   Surface SEM images of GH4151 alloy fracture morphologies after tensile testing at 550 oC

(a) cleavage fracture (b) ductile tearing


图9

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图9   23~550 ℃拉伸过程中GH4151合金断裂过程示意图

Fig.9   Schematics of fracture processes of GH4151 alloys during tensile testing at 23-550 oC (σ—tensile stress)


混合断裂中的沿晶断裂通常与晶界处硬质析出相有关[28]。如图3和10所示，在断口横截面和纵截面均可观察到富Nb、Ti的MC碳化物，并且纵截面上可观察到MC碳化物破碎产生的孔洞。如图10b和c所示，MC碳化物周围存在高密度位错。在变形过程中，位错运动被MC界面阻碍，随着应变的累积，MC碳化物周围形成位错塞积，导致MC碳化物界面处产生局部应力集中，进而引发沿晶断裂。

图10

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图10   550 ℃拉伸后GH4151合金断口纵截面的SEM像、EDS面分布图和TEM像

(b) BF image (c) DF image

Fig.10   Longitudinal SEM image and EDS elemental mappings (a) and TEM images (b, c) of the fracture morphology of GH4151 alloy after tensile testing at 550 oC


3.2 650~800 ℃拉伸过程中合金的断裂机制

在650~800 ℃温度范围内，GH4151合金塑性下降。与23~550 ℃拉伸样品不同，650~800 ℃拉伸样品中裂纹以沿晶断裂方式萌生，以混合断裂方式扩展。图11为800 ℃拉伸后GH4151合金断口纵截面的EBSD分析，其中箭头所示为一次γ′相。局部取向差(KAM)图和几何必需位错(GND)密度图显示，位错主要在晶界处累积，尤其是一次γ'相附近的晶界。如2.3节所述，650~800 ℃拉伸过程中，变形同时发生在γ基体和γ′相内，与23~550 ℃拉伸相比，更多位错可通过剪切二次γ'相的方式运动至晶界。晶界处位错运动是协调相邻两晶粒变形的一种机制[29,30]，而一次γ′相阻碍了位错运动，导致晶界局部应力集中，这些具有较高位错密度的晶界有可能是裂纹萌生的优先位置。

图11

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图11   800 ℃拉伸后GH4151合金断口的EBSD分析

Fig.11   Inverse pole figure (IPF) (a), kernel average misorientation (KAM) map (b), and geometrically necessary dislocation (GND) density map (c) of GH4151 alloy after tensile fracture at 800 oC (Black arrows represent primary γ′ phases, the same in Figs.12 and 15)


图12为800 ℃拉伸后GH4151合金裂纹尖端的SEM像和EPMA元素面分布图。可以看出，裂纹内部聚集大量O元素。因此，沿晶断裂不仅由局部应力集中所致，还可能涉及氧化损伤导致的晶界脆化。尽管本工作中拉伸速率较快，但合金中仍可观察到氧化损伤晶界现象。事实上，氧化损伤导致晶界开裂的过程极为迅速，这一点在Cao等[31] (拉伸速率为1 × 10-3 s-1)和Németh等[13] (拉伸速率为1 × 10-4 s-1)的工作中已得到证实，Molins等[32]更是认为材料脆化所需时间小于20 s。图13为真空条件和大气气氛下800 ℃拉伸时GH4151合金的应力-应变曲线和真空拉伸断口SEM像。与大气气氛拉伸样品相比，真空拉伸样品的断后伸长率大幅度提高，且断口呈现出与23~550 ℃拉伸断口相同的韧性断裂特征，即以穿晶断裂为主的混合断裂，且无明显脆性沿晶断裂区域，这充分证实了氧化损伤对650~800 ℃拉伸样品晶界的脆化作用。研究[13,31]表明，这种沿晶断裂的产生与应力辅助O原子扩散或动态脆化有关。晶界处大量缺陷聚集为O原子提供了通道[33,34]，使高能晶界成为破坏性原子的快速扩散路径，加速了650~800 ℃下O原子向晶界弹性应力场或缺陷应力场的富集，导致裂纹尖端晶界局部内聚力损失，发生动态晶界脱黏，进而引起裂纹萌生和扩展。新扩展的裂纹尖端又会重新被O原子吸附，如此往复，形成沿晶断裂。因此，可推测GH4151合金的沿晶断裂是650~800 ℃下高应力与O原子共同作用的结果。

图12

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图12   800 ℃拉伸后GH4151合金裂纹尖端处SEM像和EPMA元素面分布图

Fig.12   SEM image (a) and EPMA elemental mappings (b-i) of crack tip in GH4151 alloy after tensile testing at 800 oC (Lv. represents degree of element concentration)

(b) Al (c) Nb (d) Ti (e) O (f) Ni (g) Mo (h) Co (i) Cr


图13

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图13   在真空和大气气氛下800 ℃拉伸时GH4151合金的应力-应变曲线和真空拉伸断口SEM像

Fig.13   Tensile stress-strain curves of GH4151 alloy after tensile fracture at 800 oC in vacuum and atmospheric atmospheres (Inset in Fig.13a is photo of samples after tensile testing) (a) and SEM images of fracture of GH4151 alloy at 800 oC in vacuum (b, c)


图14为650~800 ℃拉伸过程中GH4151合金断裂过程示意图。此阶段位错可通过剪切方式切过γ'相，导致晶界位错累积程度增加，而位错在γ/γ'界面的累积程度减小。在γ/γ'界面位错累积未达到临界值之前合金便发生沿晶断裂，合金断后伸长率下降。随着应变增加，裂纹扩展速率增加，减少发生动态脆化的时间，导致动态脆化不完全[35,36]。当γ/γ'界面位错累积达到临界值时，断裂方式转变为以穿晶断裂为主的混合断裂，这种从沿晶断裂到混合断裂的转变被认为是氧化损伤与机械损伤之间竞争的结果。

图14

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图14   650~800 ℃拉伸过程中GH4151合金断裂过程示意图

Fig.14   Schematics of fracture processes of GH4151 alloy during tensile testing at 650-800 oC


3.3 950 ℃拉伸时合金的断裂机制

950 ℃拉伸后样品的塑性增加，裂纹的萌生和扩展均为氧化损伤晶界控制的沿晶断裂。如图6f所示，二次γ'相形貌转变为近似球形，γ/γ'界面共格应变强化效果减弱，对位错的阻碍作用降低。整个拉伸过程未发生机械损伤转变，说明950 ℃拉伸样品中位错在γ/γ'界面累积程度不足，不能引起γ/γ′界面微孔洞合并。如2.3节所述，在950 ℃变形时位错更容易运动至晶界处。图15为950 ℃拉伸样品断口附近显微组织的EBSD分析。与800 ℃拉伸样品相比，950 ℃拉伸样品中晶界或一次γ′相处的位错密度显著提升，表明此阶段位错更易在晶界处累积。晶界结合力与温度成反比，随实验温度升高，O原子扩散速率增加，进一步弱化晶界，导致950 ℃拉伸样品更容易产生沿晶断裂。图16为950 ℃拉伸过程中GH4151合金断裂过程示意图。在位错累积阶段，位错在γ/γ′界面累积程度急剧减小，在晶界处累积程度迅速增加，在γ/γ′界面位错累积达到临界值之前材料便完成断裂。此外，裂纹扩展速率随裂纹尖端所受应力增加而增加[37,38]，由于950 ℃拉伸后合金强度下降，外界施加应力降低，裂纹以较慢的速率扩展，导致断后伸长率增大。

图15

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图15   950 ℃拉伸后GH4151合金断口的EBSD分析

Fig.15   IPF (a), KAM map (b), and GND density map (c) of GH4151 alloy after tensile fracture at 950 oC


图16

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图16   950 ℃拉伸过程中GH4151合金断裂过程示意图

Fig.16   Schematics of the fracture processes of GH4151 alloy during tensile testing at 950 oC


4 结论

(1) 随着温度升高，GH4151合金的塑性呈现出先减小后增大的趋势，23~550 ℃拉伸样品的断裂方式为以穿晶断裂为主的混合断裂，变形主要发生在γ通道，位错在γ/γ'界面大量堆积，导致γ/γ'界面间隙撕裂形成微孔洞，从而引起穿晶断裂。混合断裂中的沿晶断裂是由于MC碳化物界面处应力集中产生的孔洞所导致。

(2) 拉伸温度为650~800 ℃时，裂纹以沿晶断裂方式萌生，随后以混合断裂方式扩展，变形同时发生在γ通道和γ′相，晶界位错塞积加速了高应力条件下O原子向晶界弹性应力场或缺陷处富集，导致晶界发生动态脆化，引起沿晶断裂，造成合金塑性下降。随着应变增加，裂纹扩展速率增大，减少了O原子导致晶界动态脆化所需时间，当γ/γ'界面位错累积达到临界值时，裂纹扩展转变为以穿晶断裂为主的混合断裂。

(3) 950 ℃拉伸时，裂纹以沿晶断裂方式萌生和扩展，γ'相形貌转变为近似球形，对位错的阻碍作用降低，不能引起γ/γ'界面微孔洞的合并，因此950 ℃拉伸样品未发生穿晶断裂。950 ℃拉伸后，合金强度下降，外界施加应力降低，裂纹以较慢的速率扩展，因此合金的塑性增加。

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采用冲击实验研究了不同温度固溶处理及固溶和双级时效处理对GH4151合金冲击性能的影响,采用光学显微镜、扫描电镜和透射电镜等观察了合金经不同工艺处理后的显微组织、析出相和断口形貌。结果表明:对于仅固溶处理的合金,当固溶温度低于1140℃时,其冲击韧性值的变化不大,固溶温度高于1140℃时,冲击韧性值随固溶温度的升高而增加,断裂过程以裂纹的稳定扩展区为主,断口呈韧性断裂。对于固溶时效处理的合金,当固溶温度低于1160℃时,冲击韧性值的变化也比较稳定,固溶温度高于1160℃时,冲击韧性值随固溶温度的升高而降低,断裂过程不存在裂纹的稳定扩展区,断口呈典型的沿晶脆性断裂。这主要是因为当固溶温度低于1140℃时,合金的晶粒尺寸变化不大,而在1140℃以上时晶粒尺寸明显增大;双级时效后,当固溶温度高于1160℃时,晶界上连续或不连续分布的第二相显著降低了合金的冲击韧性。

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