Numerical simulation and experimental evaluation of the corrosion susceptibility of degradable magnesium-based bone implants
-
摘要:
镁基合金作为"第三代生物医用材料"，以其优异的生物相容性和可降解性吸引了众多学者的关注。在传统医疗骨植入器械难以降解的背景下，展现出了其独特的潜力。然而，由于植入后在人体体液环境的应力和腐蚀作用下腐蚀损伤和退化速率的未知性，难以预测其在临床应用中的性能。因此，为防止因退化过快而导致过早地断裂失效，本研究建立了一种应用于金属腐蚀性能预测的数值模型，该模型被用于研究镁基合金的腐蚀行为预测。通过体外腐蚀试验评估其对不同程度应力腐蚀的敏感性并校准模型参数。为测试模型的准确性，对骨植入物的腐蚀行为进行了预测。结果表明，该模型可以可靠地预测各种应力条件下植入物的腐蚀损伤、降解率和相关的机械性能退化。因此，应力腐蚀模型作为一种数值模拟工具，具备准确预测腐蚀行为的同时优化植入物的退化速率的潜力。此外，提出的模型程序和方法适用于不同合金成分、多种应用场景接骨板的应力腐蚀预测，有助于实现植入材料退化速率的精确调节。
Abstract:Magnesium-based alloys, as "third-generation biomedical materials", have attracted the attention of many scholars for their excellent biocompatibility and degradability. In the context of traditional medical bone implant devices, which are difficult to degrade, it shows its unique potential. However, it is difficult to predict its performance in clinical applications due to the unknown rate of corrosion damage and degradation under the stress and corrosive effects of the human body fluid environment after implantation. Therefore, in order to prevent premature fracture failure due to excessive degradation, a numerical model applied to the prediction of corrosion performance of metals was developed in this study to investigate the in vitro corrosion behaviour of AZ31B magnesium alloy splints, and the stress-corrosion susceptibility of the splints was assessed by in vitro corrosion tests, and several sets of tests were conducted to obtain data to calibrate the model parameters. The results show that the established model can accurately predict the corrosion damage, degradation rate, and mechanical property damage of implants under different stress loads. Therefore, the stress corrosion model has the potential as a numerical simulation tool to accurately predict corrosion behaviour while optimising the degradation rate of implants. In addition, the procedures and methods of the proposed model are applicable to the stress corrosion prediction of splints with different alloy compositions and mul- tiple application scenarios, which can help to achieve an accurate regulation of the degradation rate of implant materials.
-
连续碳化硅纤维增强碳化硅复合材料(SiCf/SiC)具有轻质、耐高温、抗氧化、抗辐照、热稳定性高等优点，在航空发动机与核反应堆领域具有广泛的应用前景[1-3]。SiCf/SiC及其他SiC纤维增强的复合材料在中温氧化环境下出现强度降低及蠕变断裂时间下降等反常现象[4-6]，即中温脆化。Forio等研究了600~800℃氧化环境下单束Nicalon纤维的蠕变断裂行为，结果表明在较低应力下即发生中温脆化现象[7]，认为这种现象与晶界处游离碳的氧化有关，裂纹尖端的游离碳被氧化后导致裂纹长度发展至临界尺寸，从而导致纤维失效。在800℃的空气条件下，Nicalon/PyC/SiC会因其纤维表面氧化而出现明显的强度退化现象，断裂应力仅为拉伸强度的1/3[8]，Lara-Curzio[9]认为其持久强度σr与蠕变断裂时间t存在的关系。Morscher[10]研究了Hi-Nicalon/BN/SiC在815℃空气下的高温持久性能。与同条件下的纤维束相比，材料的承载能力显著降低，认为这种性能退化与BN界面氧化导致纤维/基体界面的强结合有关。
上述研究结果表明：中温脆化现象普遍存在于SiCf/SiC复合材料体系。本课题组的前期相关工作表明：国产二代2D-SiCf/SiC复合材料的中温脆化现象与蠕变应力大小密切相关，当蠕变应力高于比例极限应力σPLS时，出现中温脆化现象，表现出与低应力下不同的断裂机制[11]。为此，本研究采用微区结构分析与纤维推入结合的方法对中温蠕变前后的国产二代SiCf/SiC复合材料进行分析，进而探讨纤维/基体界面的微观结构与其力学性能对中温蠕变性能的影响，从而理解国产二代2D-SiCf/SiC复合材料的中温脆化机制。
1. 实验材料及方法
1.1 实验材料
本试验所用材料为国产二代纤维增强的2D-SiCf/SiC复合材料。使用的二代低氧高碳型Cansas 3200 SiC 纤维由福建立亚新材有限公司生产，纤维平均直径、平均拉伸强度和模量分别为13 μm、2.5 GPa、270 GPa；使用化学气相沉积 (Chemical Vapor Deposition, CVD) 在纤维表面制备BN界面，其厚度约30 nm；使用化学气相渗透 (CVI) 工艺制备SiC基体。复合材料的纤维体积分数、孔隙率和密度分别为35 vol%、12%和2.5 g·cm−3。中温空气蠕变试验在电子蠕变松弛试验机(RDL50，长春机械科学研究院)上进行，试验温度为500~
1000 ℃，蠕变应力为100~160 MPa，其中温蠕变性能见文献[11]。本文挑选各温度蠕变应力120 MPa(大于比例极限应力σPLS)的试样进行相关微观结构和界面力学性能的研究，见表1。表 1 本文研究试样的相关信息Table 1. Information about samples used in this studyConditions Rupture
time/hCreep
Rate/s−1Creep
Strain/%500℃/120 MPa 490 7.1×10−10 0.37 800℃/120 MPa 22 5.4×10−9 0.19 1000 ℃/120 MPa33 1.7×10−8 0.39 1.2 界面微观结构分析
采用聚焦离子/电子双束电镜(FIB, FEI Helios G4 CX)制备透射样品，并采用高分辨透射电子显微镜(TEM, FEI Talos F200X)观察蠕变前后纤维以及界面区域的精细微观结构和微区元素分布。
1.3 纤维推入实验
在纳米压痕测试系统(Hysitron, Ti980)上进行纤维推入试验，由此来分析国产二代2D-SiCf/SiC复合材料的界面结合情况。图1为纤维推入的典型载荷-位移曲线。
曲线分三个阶段：第一阶段为短暂的非线性阶段，对应压头和纤维的不完全接触过程；第二阶段为线性段，斜率为S0。此时纤维与压头完全接触，并在压头作用下产生弹性变形；第三阶段对应界面脱粘的过程。当达到临界载荷Pc后，裂纹在界面处产生并发生偏转，此时可根据剪滞模型，按式(1)计算界面剪切强度(IFSS)[13]，IFSS是评价界面性能的关键参数之一，表征了界面抵御剪切载荷的能力：
IFSS=S0Pc2π2r3Ef (1) 式中：r为纤维半径；Ef为纤维模量，其中原始纤维模量取270 GPa。
考虑到界面脱粘实质上是微裂纹在剪应力的作用下扩展的过程，根据式(2)可计算界面脱粘能(Gi)[14,15]，表示界面脱粘过程中的能量释放速率，也是评价界面性能的重要指标之一。
Gi=P2c8π2r3Ef (2) 当压头与纤维的直径比大于0.6时，纤维推入过程中的相关界面脱粘行为与压头无关[12]。本文所使用的Cansas3200系列纤维平均直径为13 μm，因此选取直径10 μm的平压头进行试验。试验采用的是位移控制模式，最大载荷为450 mN，加载速率为30 nm/s。纤维push-in测试位置的选取见图2，即具有良好的六方紧密堆积、界面完好无破损、周围无大孔洞。每个试样选取10根纤维进行试验，计算其平均性能作为材料的界面力学性能。
![]()
图 2 国产二代2D-SiCf/SiC复合材料纤维 push-in测试的位置选择Figure 2. SEM images of positions on the domestic 2nd-2D-SiCf/SiC for fiber push-in testing2. 实验结果
2.1 纤维微观结构
图3显示了不同蠕变条件下国产二代2D-SiCf/SiC复合材料中SiC纤维的HRTEM图。国产二代SiC纤维属于低氧高碳型纤维，因此在图3(a)原始试样的高分辨像(HRTEM)中可以观察到一定量的游离碳。接着对图3(b)-图3(d)中纤维蠕变后的晶粒尺寸进行统计，发现：蠕变后纤维的晶粒尺寸与原始条件下纤维的晶粒尺寸相差不大，这说明国产二代纤维在中温条件下结构稳定。研究表明，Nicalon纤维在高于
1100 ℃时才会产生明显的蠕变损伤[17]，而对于国产二代纤维来说，当温度高于1300 ℃时，才会产生明显的晶粒粗化[18,19]。因此可以认为国产二代2D-SiCf/SiC复合材料的中温脆化现象与纤维的蠕变无关。![]()
图 3 国产二代2D-SiCf/SiC复合材料中纤维的微观结构演变：(a) 原始试样；(b) 500℃/120 MPa；(c) 800℃/120 MPa；(d)1000 ℃/120 MPaFigure 3. Microstructure evolution of fibers in the domestic 2nd-2D-SiCf/SiC：(a) As-received；(b) 500℃/120 MPa； c) 800℃/120 MPa；(d)1000 ℃/120 MP2.2 界面微结构演变
图4为蠕变前后国产二代2D-SiCf/SiC复合材料界面区域的明场像和元素分布，可以看出原始试样(图4(a))的基体与界面间存在约40 nm厚的富碳层，这可能与基体制备工艺有关，BN界面中的氧含量明显高于基体和纤维。根据N元素的分布确定BN界面位置，其厚度约为30 nm。500℃/120 MPa蠕变后(图4(b))，纤维一侧出现C元素的轻微富集，并形成孔洞。800℃/120 MPa蠕变后(图4(c))，对BN界面/SiC基体和BN界面/SiC纤维区域进行EDS分析，可知这两个区域中的硅氧原子百分比分别为27.4∶53.9和31.9∶64.2，接近SiO2的原子比[20]。靠近BN界面/SiC基体侧界面的氧化较轻，该处的SiO2可能来自基体氧化，BN界面与SiC纤维中残余氧的富集使靠近BN界面/SiC纤维侧区域氧化较为严重，并形成50 nm厚的SiO2层。国产二代2D- SiCf/SiC在
1000 ℃/120 MPa条件蠕变后(图4(d))，氧元素偏聚在纤维一侧，这说明界面脱粘主要发生在BN界面/SiC纤维一侧，从而导致氧化性介质侵入界面孔隙发生化学反应，生成硅酸盐玻璃等氧化产物，从而降低纤维/基体界面结合强度，进一步导致国产二代2D-SiCf/SiC发生失效断裂。![]()
图 4 国产二代2D-SiCf/SiC复合材料蠕变前后界面区域的明场TEM像和元素分布(Si、N、O、C)：(a)原始试样；(b) 500℃/120 MPa;(c) 800℃/120 MPa；(d)1000 ℃/120 MPaFigure 4. Bright fields TEM images and element distribution of interfacial regions of the domestic 2nd-2D-SiCf/SiC before and after intermediate temperature creep (Si、N、O、C): (a) As-received; (b) 500℃/120 MPa; (c) 800℃/120 MPa; (d)1000 ℃/120 MPa为进一步探究中温蠕变对界面微观结构的影响，对典型试样的界面区域高分辨像(HRTEM)进行分析，如图5所示。图5(a)说明原始试样的BN界面区呈现典型的涡轮层状结构，这种结构表明界面沉积温度大致为900℃，低温沉积工艺对纤维布的损伤小，但得到的界面结晶度较低，抗氧化性能较差。图5(b)为500℃/120 MPa下的界面HRTEM，可见国产二代SiCf/SiC复合材料蠕变后的BN界面区域可分为两层，靠近纤维和基体区域为非晶的过渡层，而界面中心区域则为原子层相互交织的t-BN层。中心区域局部原子层排列规则，其层间距离为0.34 nm，接近h-BN的(0002)晶面间距0.33 nm，说明氧化使得BN界面原子排列更加规则，结晶度提高，具有更高的各向异性，层内结合紧密，呈现出BN纳米晶的形貌结构。图5(c)中间区域由部分BN纳米晶与非晶区域构成，界面/纤维与界面/基体结合处为非晶区，氧化程度高，结合紧密。温度进一步升高至
1000 ℃(图5(d))，BN纳米晶的含量进一步提高。![]()
图 5 国产二代SiCf/SiC界面区域的微观结构演变：(a) 原始试样；(b) 500℃/120 MPa；(c) 800℃/120 MPa；(d)1000 ℃/120 MPaFigure 5. Microstructure-evolution of interfacial region in domestic 2nd-generation SiCf/SiC: (a) As-received; (b) 500℃/120 MPa;(c) 800℃/120 MPa; (d)1000 ℃/120 MPa2.3 界面力学性能
图6为国产二代 2D-SiCf/SiC 复合材料在不同蠕变条件下纤维推入试验的典型载荷-位移曲线。500℃/120 MPa条件的推入曲线中各阶段的过渡较为平滑，无明显的界面脱粘特征，与C/SiC及碳纤维树脂基复合材料的特征相似[15]。800℃/120 MPa和
1000 ℃/120 MPa的条件下均可直接观察到明显的脱粘阶段，在脱粘点出现明显的偏折，曲线呈台阶状。![]()
图 6 不同条件下的国产二代2D-SiCf/SiC复合材料的纤维推入试验载荷-位移曲线Figure 6. Load-displacement curves obtained of the push-in tests of the domestic 2 nd-2D-SiCf/SiC.本文采用最小二乘法拟合得到线性阶段的斜率S0，并采用不同的方法对临界载荷Pc进行测定。对于脱黏后出现明显偏折的曲线，可根据曲线拐点的纵坐标直接作为Pc。对于500℃/120 MPa条件下过渡较为平滑的情况，可根据Rodrí-guez提出的“位移截距法”确定Pc[13]。先采用最小二乘法拟合出线性阶段，将得到的直线延长并与x轴相交，将该直线分别向平移至2%和10%截距处，分别与推入曲线相交于A、B两点，过该两点的直线与线性段延长线的交点即为Pc。根据上述方法和1.2节的公式所确定的界面力学性能如表2所示。
表 2 基于纤维推入获得的国产二代2D-SiCf/SiC复合材料的界面力学性能Table 2. The interfacial mechanical properties of domestic 2nd-2D-SiCf/SiC generated from push-in testConditions Pc/mN S0/(N·mm−1) IFSS/MPa Gi/(J·m−2) As-received 168±35 614±61 97±27 5.0±2.1 500℃/120 MPa 226±24 829±51 128±9 8.8±1.8 800℃/120 MPa 350±38 962±60 231±27 21.2±4.5 1000 ℃/120 MPa300±52 848±73 180±44 15.8±5.3 Notes：Pc is the critical load before interface debonding; S0 is the slope of the second stage; IFSS is the interfacial shear strength; Gi is the interphase debond energy. 在500~
1000 ℃蠕变过程中，氧化介质以基体裂纹、孔隙等作为扩散通道，与BN界面或纤维直接接触，此时SiC基体的氧化速率较低，形成的SiO2不能充分填充裂纹，界面和纤维的氧化损伤较为严重[16]。BN界面在氧化环境下形成的B2O3增大了SiC纤维或基体的溶解度，生成的SiO2部分甚至完全替代了原有的BN界面，形成界面强结合。从表中可以看出，蠕变后界面的结合程度均增大，试样在800℃/120 MPa蠕变后的界面结合最强，IFSS和Gi分别达到231 MPa和21.2 J·m−2，此时裂纹偏转等增韧机制无法发挥作用，当基体裂纹扩展至界面时，可直接贯穿纤维，导致中温脆化。当温度为500℃时，界面中纤维侧的富碳相层发生氧化反应形成空洞，BN界面发生轻微氧化，并生成BN纳米晶，材料易产生内部脱黏。随着温度升高至1000 ℃，纤维push-in曲线呈现弱平台，为典型的弱界面特征。3. 分析与讨论
国产二代2D-SiCf/SiC复合材料的界面力学性能、蠕变断裂时间与温度的关系如图7所示。从中可以发现，随着温度的升高，界面剪切强度IFSS以及界面脱黏能Gi呈现出先上升后下降的趋势。国产二代2D-SiCf/SiC复合材料室温下的σPLS约为115 MPa，因此中温条件下，蠕变应力为120 MPa，足以使基体产生一定程度的开裂，为氧化介质侵入材料内部提供了通道。500℃时，BN界面氧化较为缓慢，界面仍能维持一定的良好结合状态。当温度升高至800℃时，部分BN界面因基体裂纹直接暴露在氧化环境中，高氧分压下BN界面和SiC基体同时氧化，形成了硼硅酸盐玻璃(SiO2·B2O3)[10,21]，B2O3随空气中水分挥发后，留下一定厚度的SiO2，形成界面/基体侧的强结合；此外，由于国产二代2D-SiCf/SiC复合材料靠近纤维侧BN界面的含氧量为14% atm，较高的含氧量导致BN界面与SiC纤维在中温下自发氧化，在面/纤维一侧也形成一层SiO2，形成界面/纤维侧的强粘结，此时界面已完全失去相应的增韧效果。当温度升高至
1000 ℃时，SiC基体的氧化速率显著提高[22]，基体裂纹表面可生成大量的熔融态SiO2，熔融流动的SiO2能起到封填国产二代2D-SiCf/SiC的裂纹和孔隙等损伤或缺陷的作用，从而降低空气中的水分和氧气进入材料内部的速度，延缓了界面的氧化过程，因此，国产二代2D-SiCf/SiC在1000 ℃时的IFSS与Gi较低。在800℃/120 MPa条件下，国产二代2D-SiCf/SiC复合材料的IFSS和Gi最高，此时界面的氧化损伤也是最为严重的，较大的IFSS和Gi均说明该条件下界面处不易通过脱黏以及裂纹偏转等形式耗散断裂能，这也是强基体/界面结合的特征。界面结合的程度与蠕变断裂时间呈现较为明显的反比关系。强基体/界面结合发生后，载荷传递模式就会发生改变。当强基体/界面结合区的某处0°纤维断裂后，剩余载荷不再由整体纤维束均匀承担，而是由强基体/界面结合区的邻近纤维束承担，这大大加速了强基体/界面结合区的纤维束的失效。因此，国产二代2D-SiCf/SiC复合材料在800℃/120 MPa时的蠕变断裂时间较短。
![]()
图 7 国产二代2D-SiCf/SiC复合材料的界面力学性能、蠕变断裂时间与温度的关系：(a) IFSS；(b) GiFigure 7. Mechnical properties of fiber/matrix interphase and creep rupture time of domestic 2 nd-2D-SiCf/SiC at different temperatures: (a) IFSS; (b) Gi4. 结 论
(1) 国产二代2D-SiCf/SiC复合材料在中温高于σPLS的条件下出现中温脆化现象，与界面的氧有关。当基体开裂后，氧化介质通过基体裂纹输送至材料内部，导致BN界面氧化消耗后被SiO2填充，导致界面的强结合。
(2) 国产二代2D-SiCf/SiC复合材料的中温持久寿命与界面性能密切相关，当界面失去增韧效果后，裂纹迅速扩展并贯穿承载纤维，导致蠕变断裂时间较短。这也说明提高界面的抗氧化能力或增强SiCf/SiC复合材料的比例极限应力可有效提高材料的中温蠕变断裂时间。
-
![]()
图 1 腐蚀模拟算法流程图
Figure 1. Flowchart of corrosion simulation algorithm
D, Dc, Dh, Dc-1 and Dh-1 are the damage field, pitting damage, stress corrosion damage, pitting damage at the last time step, and stress corrosion damage at the last time step, respectively; δU and δSC are the critical thickness of the corrosion product film and the width of the corrosion cracks; Le and KU are the length of the characteristic finite element and material kinetic related parameters of the corrosion process; S and R are corrosion process and pH related parameters; Δt and α are numerically calculated time increments and time amplification constants; λe,t is the dimensionless pitting corrosion parameter of cell e at time t; σ*e is the maximum effective stress (i.e. Mises stress) of the load
![]()
图 2 (a)用于试验测试的样品；(b)(a)中测试样品的加载模型；(c)接骨板的有限元模型；(d)接骨板的结构，接骨板中螺纹孔的放大视图
Figure 2. (a) Sample used for experimental testing; (b) Loading model of the test sample in (a); (c) Finite element model of the splint; (d) Structure of the splint, enlarged view of the threaded holes in the splint
![]()
图 4 恒定载荷浸泡腐蚀设备示意图
Figure 4. Schematic diagram of constant load immersion corrosion equipment
![]()
图 5 镁合金在静载荷条件下的腐蚀过程示意图
Figure 5. Schematic diagram of the corrosion process of magnesium alloy under loaded conditions
![]()
图 6 腐蚀过程中AZ31B镁合金的质量损失率与析氢增量
Figure 6. Mass loss rate and hydrogen precipitation increment of AZ31B magnesium alloys during corrosion
![]()
图 7 AZ31B镁合金接骨板样品损伤程度随模拟时间的变化图，左下角插入的图片显示了螺纹孔附近腐蚀程度随时间的演化(点蚀与应力腐蚀协同作用)
Figure 7. Plot of the extent of damage to a sample of AZ31B magnesium alloy splints over simulation time, with the image inserted in the lower left corner showing the evolution of the extent of corrosion near the threaded holes over time (pitting corrosion in synergy with stress corrosion)
![]()
图 8 模型预测AZ31B镁合金的屈服强度在不同时间节点的损伤程度以及损伤后的屈服强度
Figure 8. Model prediction of yield strength of AZ31B magnesium alloy at different time points of damage and yield strength after damage
![]()
图 9 相同时间下，不同载荷AZ31B镁合金接骨板的损伤，插入的图片展示了关键腐蚀位置
Figure 9. Damage to AZ31B magnesium alloy splints with different loads at the same time, inserted image enlarged Critical corrosion locations
![]()
图 10 AZ31B镁合金接骨板在不同静载荷条件下(0、0.5、0.65、0.8 MPa)浸泡30 h后的宏观腐蚀形貌与模型预测结果
Figure 10. Macroscopic corrosion morphology and model prediction of AZ31B magnesium alloy splints after 30 h immersion under different static load conditions (0, 0.5, 0.65, 0.8 MPa)
![]()
图 11 模型预测AZ31B镁合金的质量损失与体外试验测得质量损失的对比图。条带表示试验中质量损失的最大值和最小值的范围，而误差线表示模型预测的质量损失的平均值±标准差
Figure 11. Plot of model-predicted mass loss of AZ31B magnesium alloy against mass loss measured in in vitro tests. The bars indicate the range of maximum and minimum values of mass loss in the test, while the error line indicates the mean ± standard deviation of the model-predicted mass loss
表 1 AZ31B的化学成分(wt%)
Table 1 Chemical composition of AZ31B (wt%)
Al Zn Mn Si Ca Fe Ni Cu Other Mg 2.74 1.27 0.361 0.0161 0.0078 0.0152 0.0016 0.0008 0.08 Bal 
下载: 导出CSV
表 2 Hank's溶液的化学成分(g·L−1)
Table 2 Chemical composition of Hank's solution (g·L−1)
NaCl KCl KH2PO4 MgSO4•7 H2O NaHCO3 CaCl2 Na2HPO4•H2O Glucose 8.00 0.40 0.06 0.20 0.35 0.14 0.06 1.00
下载: 导出CSV
表 3 AZ31B合金的腐蚀性能参数
Table 3 Corrosion performance parameters of AZ31B alloy
Load/MPa δU/mm δSC/mm KU S R Υ ψ α 0.5/0.65/0.8 0.17 0.07 0.026 0.02 3.2 0.6 5 103 Notes: δU and δSC are the critical thickness of the corrosion product film and the width of the corrosion cracks; KU is a parameter related to the corrosion process and material dynamics; S and R are corrosion process and pH related parameters; Υ and ψ are the shape and scaling parameters of the Weibull distribution, respectively; α is the time amplification constant in the numerical calculation.
下载: 导出CSV
-
[1] STAIGER M P, PIETAK A M, HUADMAI J, et al. Magnesium and its alloys as orthopedic biomaterials: a review[J]. Biomaterials, 2006, 27(9): 1728-1734. DOI: 10.1016/j.biomaterials.2005.10.003
[2] GEETHA M, SINGH A K, ASOKAMANI R, et al. Ti based biomaterials, the ultimate choice for orthopaedic implants–a review[J]. Progress in materials science, 2009, 54(3): 397-425. DOI: 10.1016/j.pmatsci.2008.06.004
[3] HUSE E C. A new ligature[J]. The Chicago Medical Journal and Examiner, 1878, 37(2): 171.
[4] WITTE F. The history of biodegradable magnesium implants: a review[J]. Acta biomaterialia, 2010, 6(5): 1680-1692. DOI: 10.1016/j.actbio.2010.02.028
[5] LAMBOTTE A. L’utilisation du magnesium comme materiel perdu dans l’osteosynthèse[J]. Bull Mem Soc Nat Chir, 1932, 28(3): 1325-1334.
[6] MCBRIDE E D. Absorbable metal in bone surgery: a further report on the use of magnesium alloys[J]. Journal of the American Medical Association, 1938, 111(27): 2464-2467. DOI: 10.1001/jama.1938.02790530018007
[7] HE M, CHEN L, YIN M, et al. Review on magnesium and magnesium-based alloys as biomaterials for bone immobilization[J]. Journal of Materials Research and Technology, 2023, 23: 4396-4419. DOI: 10.1016/j.jmrt.2023.02.037
[8] XIN Y, HU T, CHU P K. In vitro studies of biomedical magnesium alloys in a simulated physiological environment: a review[J]. Acta biomaterialia, 2011, 7(4): 1452-1459. DOI: 10.1016/j.actbio.2010.12.004
[9] ZBERG B, UGGOWITZER P J, LöFFLER J F. MgZnCa glasses without clinically observable hydrogen evolution for biodegradable implants[J]. Nature materials, 2009, 8(11): 887-891. DOI: 10.1038/nmat2542
[10] LI Z, GU X, LOU S, et al. The development of binary Mg–Ca alloys for use as biodegradable materials within bone[J]. Biomaterials, 2008, 29(10): 1329-1344. DOI: 10.1016/j.biomaterials.2007.12.021
[11] CHAYA A, YOSHIZAWA S, VERDELIS K, et al. In vivo study of magnesium plate and screw degradation and bone fracture healing[J]. Acta biomaterialia, 2015, 18: 262-269. DOI: 10.1016/j.actbio.2015.02.010
[12] WITTE F, KAESE V, Haferkamp H, et al. In vivo corrosion of four magnesium alloys and the associated bone response[J]. Biomaterials, 2005, 26(17): 3557-3563. DOI: 10.1016/j.biomaterials.2004.09.049
[13] NAUJOKAT H, SEITZ J M, ACIL Y, et al. Osteosynthesis of a cranio-osteoplasty with a biodegradable magnesium plate system in miniature pigs[J]. Acta biomaterialia, 2017, 62: 434-445. DOI: 10.1016/j.actbio.2017.08.031
[14] 丁琳琳, 张宜澜, 杨茂进, 等. 聚左旋乳酸可吸收接骨板在下颌骨骨折内固定术中的应用价值[J]. 创伤外科杂志, 2020, 22(4): 267-270. DING Linlin, ZHANG Yilan, YANG Maojin, et al. Application value of poly-L-lactic acid resorbable bone graft in internal fixation of mandibular fracture[J]. Journal of Trauma Surgery, 2020, 22(4): 267-270(in Chinese).
[15] SONG G L, ATRENS A. Corrosion mechanisms of magnesium alloys[J]. Advanced engineering materials, 1999, 1(1): 11-33. DOI: 10.1002/(SICI)1527-2648(199909)1:1<11::AID-ADEM11>3.0.CO;2-N
[16] SONG G, ATRENS A. Recent insights into the mechanism of magnesium corrosion and research suggestions[J]. Advanced Engineering Materials, 2007, 9(3): 177-183. DOI: 10.1002/adem.200600221
[17] ZENG R, DIETZEL W, WITTE F, et al. Progress and challenge for magnesium alloys as biomaterials[J]. Advanced engineering materials, 2008, 10(8): B3-B14.
[18] ZENG R, ZHANG J, HUANG W, et al. Review of studies on corrosion of magnesium alloys[J]. Transactions of nonferrous metals society of China, 2006, 16: s763-s771. DOI: 10.1016/S1003-6326(06)60297-5
[19] DAS P, KUMAR T S S, SAHU K K, et al. Corrosion, stress corrosion cracking and corrosion fatigue behavior of magnesium alloy bioimplants[J]. Corrosion Reviews, 2022, 40(4): 289-333. DOI: 10.1515/corrrev-2021-0088
[20] TROITSKII V V, TSITRIN D N. The resorbing metallic alloy ‘Osteosinthezit’as material for fastening broken bone[J]. Khirurgiia, 1944, 8(1): 41-44.
[21] HIRAYAMA K, TODA H, FU D, et al. Damage micromechanisms of stress corrosion cracking in Al-Mg alloy with high magnesium content[J]. Corrosion Science, 2021, 184: 109343. DOI: 10.1016/j.corsci.2021.109343
[22] BARFIELD W R, COLBATH G, DESJARDINS J D, et al. The potential of magnesium alloy use in orthopaedic surgery[J]. Current Orthopaedic Practice, 2012, 23(2): 146-150. DOI: 10.1097/BCO.0b013e31824a553b
[23] LI D, ZHANG D, YUAN Q, et al. In vitro and in vivo assessment of the effect of biodegradable magnesium alloys on osteogenesis[J]. Acta Biomaterialia, 2022, 141: 454-465. DOI: 10.1016/j.actbio.2021.12.032
[24] SANCHEZ A H M, LUTHRINGER B J C, Feyerabend F, et al. Mg and Mg alloys: how comparable are in vitro and in vivo corrosion rates? A review[J]. Acta biomaterialia, 2015, 13: 16-31. DOI: 10.1016/j.actbio.2014.11.048
[25] BAINO F, YAMAGUCHI S. The use of simulated body fluid (SBF) for assessing materials bioactivity in the context of tissue engineering: review and challenges[J]. Biomimetics, 2020, 5(4): 57. DOI: 10.3390/biomimetics5040057
[26] GALVIN E, CUMMINS C, YOSHIHARA S, et al. Plastic strains during stent deployment have a critical influence on the rate of corrosion in absorbable magnesium stents[J]. Medical & Biological Engineering & Computing, 2017, 55: 1261-1275.
[27] TÖRNE K, ÖRNBERG A, WEISSENRIEDER J. Influence of strain on the corrosion of magnesium alloys and zinc in physiological environments[J]. Acta Biomaterialia, 2017, 48: 541-550. DOI: 10.1016/j.actbio.2016.10.030
[28] SNIR Y, BEN-HAMU G, ELIEZER D, et al. Effect of compression deformation on the microstructure and corrosion behavior of magnesium alloys[J]. Journal of alloys and compounds, 2012, 528: 84-90. DOI: 10.1016/j.jallcom.2012.03.010
[29] WINZER N, ATRENS A, DIETZEL W, et al. Characterisation of stress corrosion cracking (SCC) of Mg-Al alloys[J]. Materials Science and Engineering: A, 2008, 488(1-2): 339-351. DOI: 10.1016/j.msea.2007.11.064
[30] JOSHI A, DIAS G, STAIGER M P. In silico modelling of the corrosion of biodegradable magnesium-based biomaterials: modelling approaches, validation and future perspectives[J]. Biomaterials Translational, 2021, 2(3): 257.
[31] DUDDU R. Numerical modeling of corrosion pit propagation using the combined extended finite element and level set method[J]. Computational Mechanics, 2014, 54(3): 613-627. DOI: 10.1007/s00466-014-1010-8
[32] BAJGER P, ASHBOURN J M A, MANHAS V, et al. Mathematical modelling of the degradation behaviour of biodegradable metals[J]. Biomechanics and modeling in mechanobiology, 2017, 16: 227-238. DOI: 10.1007/s10237-016-0812-3
[33] SAITO K, KUNIYA J. Mechanochemical model to predict stress corrosion crack growth of stainless steel in high temperature water[J]. Corrosion Science, 2001, 43(9): 1751-1766. DOI: 10.1016/S0010-938X(00)00173-6
[34] WENMAN M R, TRETHEWEY K R, JARMAN S E, et al. A finite-element computational model of chloride-induced transgranular stress-corrosion cracking of austenitic stainless steel[J]. Acta materialia, 2008, 56(16): 4125-4136. DOI: 10.1016/j.actamat.2008.04.068
[35] QUINN C, VAN GAALEN K, MCHUGH P E, et al. An enhanced phenomenological model to predict surface-based localised corrosion of magnesium alloys for medical use[J]. Journal of the Mechanical Behavior of Biomedical Materials, 2023, 138: 105637. DOI: 10.1016/j.jmbbm.2022.105637
[36] BOLOTIN V V, SHIPKOV A A. Mechanical aspects of corrosion fatigue and stress corrosion cracking[J]. International Journal of Solids and Structures, 2001, 38(40-41): 7297-7318. DOI: 10.1016/S0020-7683(01)00002-6
[37] GARUD Y S. Quantitative evaluation of environmentally assisted cracking: a survey of developments and application of modeling concepts[J]. Journal of Pressure Vessel Technology, 1991, 113(1): 1-9. DOI: 10.1115/1.2928722
[38] ABDALLA M, JOPLIN A, ELAHINIA M, et al. Corrosion modeling of magnesium and its alloys for biomedical applications[J]. Corrosion and Materials Degradation, 2020, 1(2): 11.
[39] GASTALDI D, SASSI V, PETRINI L, et al. Continuum damage model for bioresorbable magnesium alloy devices—Application to coronary stents[J]. Journal of the mechanical behavior of biomedical materials, 2011, 4(3): 352-365. DOI: 10.1016/j.jmbbm.2010.11.003
[40] GALVIN E, O'BRIEN D, CUMMINS C, et al. A strain-mediated corrosion model for bioabsorbable metallic stents[J]. Acta Biomaterialia, 2017, 55: 505-517. DOI: 10.1016/j.actbio.2017.04.020
[41] OPPEAL A. Experimental characterisation and finite element modeling of biodegradable magnesium stents[D]. Ghent: Ghent University, 2015.
[42] DEBUSSCHERE N, SEGERS P, DUBRUEL P, et al. A computational framework to model degradation of biocorrodible metal stents using an implicit finite element solver[J]. Annals of Biomedical Engineering, 2016, 44: 382-390. DOI: 10.1007/s10439-015-1530-1
[43] GROGAN J A, O’BRIEN B J, LEEN S B, et al. A corrosion model for bioabsorbable metallic stents[J]. Acta biomaterialia, 2011, 7(9): 3523-3533. DOI: 10.1016/j.actbio.2011.05.032
[44] BARZEGARI M, MEI D, LAMAKA S V, et al. Computational modeling of degradation process of biodegradable magnesium biomaterials[J]. Corrosion Science, 2021, 190: 109674. DOI: 10.1016/j.corsci.2021.109674
[45] SHEN Z, ZHAO M, ZHOU X, et al. A numerical corrosion-fatigue model for biodegradable Mg alloy stents[J]. Acta biomaterialia, 2019, 97: 671-680. DOI: 10.1016/j.actbio.2019.08.004
[46] HAMMI Y, GULLETT P, HORSTEMEYER M F. Multiscale Concrete Modeling of Aging Degradation[R]. Mississippi State Univ. , Mississippi State, MS (United States), 2015.
[47] 国家食品药品监督管理总局. 骨接合植入物 金属接骨板: YY 0017-2016[S]. 北京: 中国标准出版社, 2016. State Administration for Market Regulation. Implants for osteosynthesis-Metallic bone plates: YY 0017-2016[S]. Beijing: China Standards Press, 2016(in Chinese).
[48] SIMULIA D S. ABAQUS Analysis manual[J]. Provid RI, 2016, 423.
[49] HANKS J H. Hanks' balanced salt solution and pH control[J]. TCA manual/Tissue Culture Association, 1975, 1(1): 3-4. DOI: 10.1007/BF00914425
[50] 中国国家标准化管理委员会. 金属和合金的腐蚀 腐蚀试样上腐蚀产物的清除: GB/T 16545-2015[S]. 北京: 中国标准出版社, 2015. Standardization Administration of the People’s Republic of China. Corrosion of metals and alloys-Removal of corrosion products from corrosion test specimens: GB/T 16545-2015[S]. Beijing: China Standards Press, 2015(in Chinese).
[51] 郑斌. 基于有限元的智能气囊小夹板治疗桡骨远端关节内骨折的生物力学研究 [D]. 江苏: 南京中医药大学, 2016. ZHENG Bin. Biomechanical study of a finite element-based smart airbag mini-splint for the treatment of intra-articular fractures of the distal radius [D]. Jiangsu: Nanjing University of Chinese Medicine, 2016(in Chinese).
[52] 黄若景. 3D打印计算机辅助设计小夹板外骨骼系统的研究 [D]. 广东: 广州中医药大学, 2015. HUANG Ruojing. 3D printing computer-aided design of a small splinted exoskeleton system [D]. Guangdong: Guangzhou University of Chinese Medicine, 2015(in Chinese).
[53] 中国国家标准化管理委员会. 金属材料 室温压缩试验方法: GB/T 7314-2017[S]. 北京: 中国标准出版社, 2017. Standardization Administration of the People’s Republic of China. Metallic materials-Compression test method at room temperature: GB/T 7314-2017[S]. Beijing: China Standards Press, 2017(in Chinese).
[54] SONG G L, UNOCIC K A. The anodic surface film and hydrogen evolution on Mg[J]. Corrosion science, 2015, 98: 758-765. DOI: 10.1016/j.corsci.2015.05.047
[55] ASCENCIO M, PEKGULERYUZ M, OMANOVIC S. An investigation of the corrosion mechanisms of WE43 Mg alloy in a modified simulated body fluid solution: the effect of electrolyte renewal[J]. Corrosion Science, 2015, 91: 297-310. DOI: 10.1016/j.corsci.2014.11.034
[56] MELANDER A, LARSSON M. The effect of stress amplitude on the cause of fatigue crack initiation in a spring steel[J]. International Journal of Fatigue, 1993, 15(2): 119-131. DOI: 10.1016/0142-1123(93)90006-C
[57] 鲁晔. 拉伸变形对AZ31镁合金应力腐蚀行为影响 [D]. 山西: 太原理工大学, 2021. LU Ye. Effect of tensile deformation on stress corrosion behaviour of AZ31 magnesium alloy [D]. Shanxi: Taiyuan University of Technology, 2021(in Chinese).
[58] MEI D, WANG C, LAMAKA S V, et al. Clarifying the influence of albumin on the initial stages of magnesium corrosion in Hank's balanced salt solution[J]. Journal of Magnesium and Alloys, 2021, 9(3): 805-817. DOI: 10.1016/j.jma.2020.07.002
[59] DA COSTA-MATTOS H S, BASTOS I N, GOMES J. A simple model for slow strain rate and constant load corrosion tests of austenitic stainless steel in acid aqueous solution containing sodium chloride[J]. Corrosion Science, 2008, 50(10): 2858-2866. DOI: 10.1016/j.corsci.2008.07.020
[60] CHOUDHARY L, RAMAN R K S. Magnesium alloys as body implants: Fracture mechanism under dynamic and static loadings in a physiological environment[J]. Acta biomaterialia, 2012, 8(2): 916-923. DOI: 10.1016/j.actbio.2011.10.031
[61] KAPPES M, IANNUZZI M, CARRANZA R M. Pre-exposure embrittlement and stress corrosion cracking of magnesium alloy AZ31B in chloride solutions[J]. Corrosion, 2014, 70(7): 667-677. DOI: 10.5006/1172
[62] JANG Y, COLLINS B, SANKAR J, et al. Effect of biologically relevant ions on the corrosion products formed on alloy AZ31B: an improved understanding of magnesium corrosion[J]. Acta biomaterialia, 2013, 9(10): 8761-8770. DOI: 10.1016/j.actbio.2013.03.026
[63] HAN L, ZHANG Z, DAI J, et al. In vitro bio-corrosion behaviors of biodegradable AZ31B magnesium alloy under static stresses of different forms and magnitudes[J]. Journal of Magnesium and Alloys, 2023, 11(3): 1043-1056. DOI: 10.1016/j.jma.2021.09.018
[64] WANG B, GAO J, WANG L, et al. Biocorrosion of coated Mg–Zn–Ca alloy under constant compressive stress close to that of human tibia[J]. Materials Letters, 2012, 70: 174-176. DOI: 10.1016/j.matlet.2011.12.009
[65] SHI Z, SONG G, ATRENS A. Corrosion resistance of anodised single-phase Mg alloys[J]. Surface and Coatings Technology, 2006, 201(1-2): 492-503. DOI: 10.1016/j.surfcoat.2005.11.081
[66] PADEKAR B S, RAJA V S, RAMAN R K S. Stress corrosion cracking of a wrought Mg–Mn alloy under plane strain and plane stress conditions[J]. Engineering Fracture Mechanics, 2013, 102: 180-193. DOI: 10.1016/j.engfracmech.2013.02.018
[67] ZHAO D, HUANG S, LU F, et al. Vascularized bone grafting fixed by biodegradable magnesium screw for treating osteonecrosis of the femoral head[J]. Biomaterials, 2016, 81: 84-92. DOI: 10.1016/j.biomaterials.2015.11.038
[68] LEE J W, HAN H S, HAN K J, et al. Long-term clinical study and multiscale analysis of in vivo biodegradation mechanism of Mg alloy[J]. Proceedings of the National Academy of Sciences, 2016, 113(3): 716-721. DOI: 10.1073/pnas.1518238113
[69] WINDHAGEN H, RADTKE K, WEIZBAUER A, et al. Biodegradable magnesium-based screw clinically equivalent to titanium screw in hallux valgus surgery: short term results of the first prospective, randomized, controlled clinical pilot study[J]. Biomedical engineering online, 2013, 12: 1-10 DOI: 10.1186/1475-925X-12-1
-
其他相关附件
-
本文图文摘要
点击下载
-
-
目的
镁基合金作为新型医用金属材料，被称作"第三代生物医用材料"，具有优异的生物相容性和可降解性。近年来，经过国内外学者的不断研究，已经作为骨科植入物被应用于临床治疗。然而，由于植入后在人体体液环境的应力和腐蚀作用下腐蚀损伤和退化速率的未知性，难以预测其在临床应用中的性能。因此，为防止因退化过快而导致过早地断裂失效，本研究建立了一种应用于金属腐蚀性能预测的数值模型，该模型被用于研究镁基合金的腐蚀行为预测。
方法本文基于连续介质损伤力学(CDM)的镁合金腐蚀模型，建立了AZ31B镁合金可降解骨科植入物的生理腐蚀预测模型。该模型损伤机制由点蚀与应力腐蚀模型组成，结合了Gastaldi等和Grogan等提出的模型，并采用Shen等提出的损伤协同效应贡献方式。值得一提的是，接骨板在人体内主要以固定作用为主，而像血管支架等器械在作用时往往活动频繁，需要考虑疲劳损伤带来的腐蚀影响，因此这里采用了更适合接骨板实际使用情况的恒定载荷应力腐蚀模型。体外校准试验采用专门设计的自制模拟生理应力腐蚀试验装置，研究了浸没在Hank's平衡盐溶液(HBSS)中的AZ31B合金接骨板受载荷条件下的腐蚀行为以及无载荷和不同外加载荷大小对镁合金生物腐蚀性能的影响。通过试验数据检验了腐蚀模型的准确性，并进一步讨论了外加载荷对接骨板的腐蚀损失、机械性能损失与退化速率的影响。
结果(1)恒定载荷下的腐蚀浸泡试验结果显示，镁合金在无应力条件下的局部腐蚀更加严重；当外载荷作用在镁合金上时，在应力(外加应力与残余应力)与电化学腐蚀的共同作用下，镁合金表面氧化膜发生破裂，部分基体直接暴露在腐蚀介质中，腐蚀坑在暴露区域形成，随着不断腐蚀，腐蚀坑的中会形成更大应力集中现象，从而导致裂纹的产生；随着应力的增大，将加快腐蚀裂纹的产生进程，并且析氢量与外加应力大小成正比。(2)腐蚀损伤预测结果显示，点蚀和应力腐蚀模型都能够在模拟过程中预测材料的腐蚀情况，腐蚀的整个过程呈现出先点蚀的现象，大约在20h时应力腐蚀才开始发挥作用，之后进入两种腐蚀机制协同作用阶段。其中，点蚀的影响导致接骨板样品出现了大面积的腐蚀情况，而应力腐蚀影响主要集中于螺纹孔区域，并且出现腐蚀裂纹。(3)机械性能损伤预测结果显示，随着腐蚀时间增长，腐蚀损伤不断的累积导致屈服强度损伤不断增加，30h时损伤程度达到80%左右，此刻腐蚀凹坑已经遍布接骨板表面，并且螺纹孔周围出现裂纹，表明材料机械性能已基本丧失。(4)不同载荷下接骨板腐蚀损伤试验与模型对比结果显示，随着载荷的增加接骨板的腐蚀损伤程度逐渐增大，0.8MPa载荷下损伤最严重，其质量损失率达到了大约40%；无载荷下的接骨板样品在浸泡30h后表面隐约可见金属光泽，并且由于无应力集中的影响，螺纹孔的受损程度与整体相差较小，结构完整性仍然保留，并且质量损失率和析氢量最小。(5)体外试验与模型预测的退化率对比结果显示，开发的点蚀与应力腐蚀模型对于外加载荷0.5MPa接骨板的质量损失预测结果略差一些，0.8MPa接骨板的预测结果相对更准确，并且从预测的趋势不难看出，负载与准确率成正比，表明模型对于预测高负载受腐蚀接骨板的退化率更加敏感。
结论(1)恒定载荷下的浸泡试验表明，随着载荷的增加接骨板的腐蚀损伤程度逐渐增大，0.8MPa载荷下损伤最严重，其质量损失率达到了大约40%。并且不同载荷下的腐蚀损伤趋势基本相同，均呈先增加后减小的趋势。此外，腐蚀过程中受载荷的接骨板试样表面会呈现出均匀分布的平行腐蚀痕迹，载荷能够加速腐蚀进程。(2)该模型能够准确预测植入物在高应力载荷下的腐蚀损伤、机械性能损伤以及退化速率，而对植入物在低应力载荷下的退化速率预测虽然有一定的差异，但仍具有一定的适用性。(3)植入物的腐蚀断裂形貌仅仅能够被大致推断。值得一提的是，腐蚀敏感性参数应被视为合金微观组织和变化腐蚀条件的函数，以确保腐蚀模型能够准确地反映试验中观察到的腐蚀形貌。这也间接表明，探索更高级的腐蚀损伤模型对于深入理解骨植入物的复杂腐蚀行为具有重要的意义。
-
近年来，医用镁合金作为“第三代生物医用材料”，因其具有优异的生物相容性和可降解性吸引了众多学者的关注，在传统医用植入物难以降解的背景下，展现了独有的潜力。然而，由于其较差的腐蚀性能，难以满足医用植入物的临床应用需求，因此，研究医用镁合金的腐蚀性能就有着举足轻重的作用。
本研究建立了一种应用于镁基合金腐蚀性能预测的数值模型，对AZ31B镁合金接骨板的腐蚀行为进行了研究，并通过体外腐蚀试验对接骨板的应力腐蚀敏感性进行了评估。该模型损伤机制由点蚀与应力腐蚀模型组成，在过往研究人员提出的腐蚀模型基础上加以改进，并通过自制的体外腐蚀试验设备进行了多组试验校准模型，使模型的预测结果更加准确，弥合了体外试验研究和临床应用之间的差距。预测结果显示，模型能够准确预测植入物在高应力载荷下的腐蚀损伤、机械性能损伤以及退化速率，而对植入物在低应力载荷下的退化速率预测也具有一定的适用性。体外腐蚀试验结果显示，随着载荷的增加接骨板的腐蚀损伤程度逐渐增大，0.8MPa载荷下损伤最严重，其质量损失率达到了大约40%。并且不同载荷下的腐蚀损伤趋势基本相同，均呈先增加后减小的趋势。
(a)腐蚀过程中的质量损失率与析氢增量；(b)模型预测的质量损失与体外试验测得质量损失的对比。条带表示试验中质量损失的最大值和最小值的范围，而误差线表示模型预测的质量损失的平均值±标准差
(a) Rate of mass loss during corrosion versus incremental hydrogen precipitation; (b) Model-predicted mass loss versus mass loss measured in in vitro tests. The bands indicate the range of maximum and minimum values of mass loss in the tests, while the error line indicates the mean ± standard deviation of the model-predicted mass loss
计量
- 文章访问数: 101
- HTML全文浏览量: 106
- PDF下载量: 2
