退火热处理对增材制造AlSi10Mg合金宏观和微观变形行为的影响
Effect of Annealing Heat Treatment on the Macroscopic and Microscopic Deformation Behavior of Additively Manufactured AlSi10Mg Alloy
通讯作者: 张星星,xxzhang@ihep.ac.cn,主要从事金属材料的大科学装置表征与计算模拟研究
责任编辑: 李海兰
收稿日期: 2024-04-23 修回日期: 2024-05-18
基金资助: |
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Corresponding authors: ZHANG Xingxing, associate professor, Tel:
Received: 2024-04-23 Revised: 2024-05-18
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作者简介 About authors
张星星,男,1984年生,副研究员,博士
增材制造AlSi10Mg合金存在显著的残余应力,对零件的形状尺寸精度、服役安全性和抗腐蚀性能等产生不利影响。在实际应用中,对残余应力敏感的应用需采用去应力退火以消除残余应力。然而,目前对增材制造AlSi10Mg合金退火后力学特性的理解还停留在宏观层面。为了更深入地揭示其微观力学行为与内在机制,本工作利用同步辐射X射线衍射技术,对合金进行原位变形研究,分析Al和Si相的晶格应变应力演化,明确各物相对合金加工硬化率的具体贡献,并对位错密度的演变进行了量化分析,从而阐明了退火热处理对增材制造AlSi10Mg合金载荷传递行为与位错行为的影响。
关键词:
The additively manufactured AlSi10Mg alloy demonstrates considerable residual stresses, adversely affecting the dimensional accuracy, operational safety, and corrosion resistance of the parts. In practical applications, stress relief annealing is necessary to eliminate residual stresses in residual stress-sensitive applications. However, the current understanding of the mechanical properties of the additively manufactured AlSi10Mg alloy after annealing is still limited to the macroscopic level. To further investigate the micromechanical behavior and intrinsic mechanisms of the alloy, this study employed synchrotron X-ray diffraction technology to conduct in situ deformation analysis. This study thoroughly examined the lattice strain and stress evolutions of the Al and Si phases and clarified the individual contribution of each phase to the strain hardening rate of the alloy. In addition, this study quantitatively assessed the evolution of dislocation density and elucidated the influences of annealing heat treatment on the load transfer and dislocation behavior of the additively manufactured AlSi10Mg alloy.
Keywords:
本文引用格式
张星星, LUTZ Andreas, 甘为民, MAAWAD Emad, KRIELE Armin.
ZHANG Xingxing, LUTZ Andreas, GAN Weimin, MAAWAD Emad, KRIELE Armin.
目前,已有较多研究探讨了热处理制度对增材制造Al-Si-Mg合金微观结构和力学性能的影响。例如,Huang等[8]发现,LPBF制备的AlSi10Mg合金在270℃去应力退火2 h后,Al-Si共晶网络部分断裂,Si相发生半球化,导致强度降低而延伸率提高。Takata等[9]研究表明,LPBF制备的AlSi10Mg合金在300℃退火2 h后,屈服强度和拉伸强度分别由279和475 MPa降为180和285 MPa,延伸率由7.5%提高至18.6%。Chen等[10]研究表明,LPBF制备的AlSi10Mg合金在280℃退火2 h后强度降低而延伸率提高,且力学性能各向异性明显降低。Zhang等[11]研究发现,LPBF制备的AlSi10Mg合金在300℃退火2 h后,Al-Si共晶网络断裂,Si相球化,导致屈服强度、拉伸强度和疲劳强度均明显降低。Tridello等[12]的研究指出,不同的退火温度会显著影响LPBF制备的AlSi10Mg合金的超高周疲劳性能,适当的退火处理(244℃退火2 h)能够保持共晶网络的完整性,从而提高合金的疲劳强度。这些研究表明了热处理制度在调控增材制造Al-Si-Mg合金性能中的重要作用。
尽管现有研究已涉及增材制造Al-Si-Mg合金的宏观力学性能和微观组织分析,但对其微观力学行为及其机制理解仍显不足。同步辐射X射线衍射(XRD)和中子衍射原位变形实验技术能够更精确地分析金属材料在变形过程中的微观力学参量变化,从而为理解材料变形行为和力学性能提供新的、深刻的见解。Xie等[13]采用中子衍射原位循环变形实验揭示了6061铝合金的变形诱导损伤各向异性,探测到了晶体取向相关的晶格应变行为。本课题组[14]前期采用同步辐射XRD原位变形实验分析了增材制造AlSi10Mg打印态合金的晶格应变、相应力和位错密度演化行为,发现Si相存在庞载荷传递效应(giant load transfer effect),Si相最大晶格应变高达约12000 × 10-6,约为铸造合金对应值的4倍,对合金力学行为有显著影响;同时还发现了多级位错强化行为。随后,Takata等[15]采用同步辐射XRD原位变形实验分析了增材制造Al-12%Si (质量分数)打印态和300℃、2 h退火态及530℃、6 h退火态合金的晶格应变、相应力和位错密度演化行为。Li等[16]也采用同步辐射XRD原位变形实验分析了增材制造AlSi10Mg打印态合金的晶格应变和相应力演化行为。目前,增材制造AlSi10Mg合金在退火后的晶格应变、应力和位错密度演化行为还需进一步研究。
本工作基于同步辐射XRD原位变形实验技术,对LPBF垂直制备的AlSi10Mg打印态和退火态合金进行了分析,探讨合金变形过程中的晶格应变与相应力,同时采用CMWP方法量化了位错密度的演化。通过对加工硬化行为进行对比分析,阐明了退火热处理对载荷传递行为与位错行为的影响,旨在为建立精确的力学性能预测模型(如晶体塑性模型)提供参考依据。
1 实验方法
1.1 材料制备
AlSi10Mg合金棒状试样由LPBF方法垂直打印(这与前期研究[14]不同(其试样为水平打印方式)),所用设备为Concept Laser M2 cusing激光系统,实验激光参数[14]为:激光功率370 W,激光直径150 μm,激光扫描速率1500 mm/s,层厚30 μm,基板未预热。垂直打印样品分为2部分,一部分保持打印态,不进行热处理,通过机械加工制成哑铃型棒状试样,试样总长60 mm,平行段长度20 mm、直径4 mm,见图1a;另一部分样品在270℃去应力退火2 h,再机械加工成哑铃状棒型试样,试样总长60 mm,平行段长度20 mm、直径5 mm,见图1b。虽然打印态试样与去应力退火试样的平行段直径有所不同,但并不影响对实验结果的分析。
图1
图1
原位变形实验试样尺寸
Fig.1
Sample dimensions for the in situ deformation experiment (unit: mm)
(a) as-printed sample
(b) stress-relieved annealed sample
1.2 微观组织观察
打印态和去应力退火态试样经机械研磨和抛光后,采用Thermo Fischer Quattro S型扫描电镜(SEM)观察试样的微观组织。为了获得最佳的物相和晶粒取向对比,使用了圆形背散射电子探测器。对于2500放大倍数,所使用的电压为4.0 kV,电子束斑尺寸6.0,光圈直径20 μm;对于15000放大倍数,所使用的电压为10.0 kV,电子束斑尺寸4.0,光圈直径20 μm。
1.3 同步辐射X射线衍射原位拉伸变形实验
在DESY PETRA III光源的P07B线站(由亥姆霍兹联合会运行)开展同步辐射XRD原位拉伸变形实验。拉伸变形过程通过位移控制,拉伸速率0.3 mm/min,相当于2.5 × 10-4 s-1的名义应变速率。试样平行段应变由Instron 2620-602型引伸计测量。同步辐射X射线能量为87.1 keV,X射线波长0.014235 nm,光束尺寸0.7 mm × 0.7 mm。面探测器型号为PerkinElmer XRD 1622。二维衍射图的180° ± 5°和270° ± 5°分区分别用于计算横向(TD)和加载方向(LD)的衍射强度-2θ曲线。
1.4 晶格应变、应力与位错密度分析
s方向(s为LD或TD)的{hkl}晶格应变
由衍射峰拟合误差所带来的晶格应变误差
式中,
采用CMWP方法分析位错密度。在CMWP方法中,物理线型函数通过尺寸宽化和微观应变宽化函数的卷积计算获得,均方应变
式中,
2 实验结果和分析讨论
2.1 宏观力学性能
图2a显示了打印态和退火态试样的宏观拉伸曲线。其中,打印态试样拉伸至断裂,断裂时工程应变为0.10;退火态试样拉伸至0.10工程应变后卸载。图2b显示了试样的加工硬化率曲线。图2表明,打印态试样的塑性变形行为可以分为3个阶段,分别记为A、B和C阶段。阶段A的加工硬化率非常高,但随应变增加而快速降低;阶段B的加工硬化率变化很小,约为6000 MPa,因此阶段B的真应力-应变曲线近似为直线;阶段C的加工硬化率随着应变增大而逐步减小,当真应变大于0.07后,加工硬化率趋于定值。退火态试样的加工硬化率随着应变增加持续平滑减小。上述结果表明,退火后,宏观力学行为的改变不仅体现在强度降低和加工硬化率降低方面,加工硬化率的演化路径形态也发生了改变。这些变化与微观结构和微观力学行为的改变有关。
图2
图2
打印态和退火态试样的宏观力学行为
Fig.2
Macroscopic behaviors of as-built and annealed samples
(a) engineering stress-strain curves
(b) hardening coefficient curves
(c) determination of uniform true strain
根据Considère颈缩判据
2.2 微观组织
图3
图3
打印态试样和退火态试样显微组织的SEM像
Fig.3
SEM images of microstructures of as-built (a-c) and annealed (d-f) samples
(a, d) microstructures at low magnification
(b, e) microstructures within the melt-pool
(c, f) microstructures on the melt-pool boundary
2.3 晶格应变与应力演化
图4
图4
打印态和退火态试样Al基体和Si相的{hkl}晶格应变演化
Fig.4
Evolutions of lattice strains of as-built (a, c) and annealed (b, d) samples (LD—loading direction, TD—transverse direction)
(a, b) Al matrix (c, d) Si phase
表1 变形过程中晶格应变的最大误差 (10-6)
Table 1
Sample | Direction | Al{111} | Al{200} | Al{220} | Al{311} | Si{111} | Si{220} | Si{311} | Si{422} |
---|---|---|---|---|---|---|---|---|---|
As-built | LD | 16.86 | 15.58 | 20.11 | 20.57 | 182.73 | 2467.52 | 209.37 | 308.45 |
TD | 49.80 | 24.88 | 15.74 | 26.65 | 141.62 | 1574.93 | 295.66 | 281.45 | |
Annealed | LD | 29.40 | 30.40 | 14.51 | 26.19 | 83.85 | 2019.93 | 105.50 | 97.65 |
TD | 51.04 | 34.36 | 14.37 | 34.36 | 95.04 | 568.67 | 76.91 | 95.85 |
(1) 阶段1为纯弹性变形阶段:在这一阶段,Al、Si两相的{hkl}晶格应变随着加载应力的增加而线性变化。
(2) 阶段2为弹塑性转变前期:此时,Al基体开始出现塑性变形,虽然程度很小,但已导致晶格应变曲线略微偏离了原有的直线轨迹。同时,由于Al基体的塑性变形,载荷传递作用使得Si相在加载方向上的晶格应变开始加速增大。
(3) 阶段3为弹塑性转变后期:在这个阶段,Al基体的各个晶面已经发生了明显的塑性变形,晶格应变曲线显著偏离直线。其中Al{111}晶面表现出加工硬化的特性,而Al{200}、{220}和{311}晶面则呈现出加工软化的特征。同时,Si相在加载方向上的各晶格应变几乎是线性且迅速地增加。
(4) 阶段4为宏观塑性变形阶段:此阶段,Al基体中的所有晶格应变均显示出加工硬化的特征,而Si相在加载方向上的各晶格应变则继续保持快速增长。
(5) 阶段5为损伤阶段:在这一阶段,Si相在加载方向上的晶格应变缓慢增加,随后开始下降。这是由于Si相的损伤导致其承载能力减弱,部分载荷因此转移到Al基体上,使得Al基体中的各晶格应变随着加载应力的增加而迅速上升。
(1) 在阶段3,退火态试样在加载方向上的Al{200}、{220}和{311}晶面虽然也表现出加工软化现象,但晶格应变的降低幅度相对较小。
(2) 在阶段4,退火态试样在加载方向上的Al{111}与{200}晶格应变最大差异约为444 × 10-6,远小于打印态试样的1047 × 10-6。这表明退火处理有助于缓解Al基体晶粒尺度的塑性变形各向异性。
(3) 值得注意的是,退火态试样在加载方向上的Si相最大晶格应变出现在{220}晶面,为7619 × 10-6。相比之下,打印态试样的最大晶格应变出现在Si的{311}晶面,高达18267 × 10-6,是退火态试样的2倍以上。这种差异可能与退火过程中Si相球化有关,退火改变了Si相的形态、尺寸和晶体取向,进而影响其微观力学行为。
(4) 退火态试样的Si相晶格应变未出现损伤阶段,这与其较小的晶格应变有关。退火导致Al基体发生回复,降低了其强度,使得流变应力偏低,不足以支撑足够强的载荷传递效应。因此,退火态试样中Si相承受的最大载荷小于打印态,未达到引发损伤所需的临界载荷。
根据晶格应变和Hook定律可以计算出Al和Si两相中不同{hkl}晶格的应力,如图5所示。图5a和b比较了Al基体中各{hkl}晶格应力与合金的宏观应力。可以看出,在弹性变形阶段,Al基体的应力与宏观应力基本一致;而进入塑性变形阶段,合金的宏观应力明显大于Al基体中的应力,这一差异主要由Si相应力所引起。图5c和d表明,在均匀变形阶段,打印态和退火态合金中Si{311}晶格应力最大值分别达到了2885和1224 MPa。前者远远超过了前期研究结果[14,21,22]所报道的数值,在前期研究中,LPBF水平打印的AlSi10Mg合金中Si{311}晶格应力的最大值约为1948 MPa (此时合金宏观真应变为0.058)[14],而LPBF垂直打印的AlSi3.5Mg2.5合金中Si{311}晶格应力的最大值约为2363 MPa (此时合金宏观真应变为0.090)[21]。通过对比分析,并结合图5c和d结果,推断Si相中的最大应力与合金塑性密切相关:合金的应变越大,Si相应力也随之增大[23]。
图5
图5
打印态退火态样品Al基体和Si相的不同{hkl}晶格的应力演化
Fig.5
Evolutions of phase stresses of as-built (a, c) and annealed (b, d) samples
(a, b) Al matrix (c, d) Si phase
考虑到AlSi10Mg合金中Si相的体积含量约为9.63%[14],可以推算出,在打印态和退火态合金中,Si相对最大载荷的贡献率分别为2885 × 9.63% ÷ 505 = 55%和1224 × 9.63% ÷ 280 = 42%,其中,505和280 MPa分别是打印态和退火态合金中Si相Si{311}晶格应力最大时所对应的合金宏观真应力。该结果揭示了Si相的庞载荷传递强化效应,即仅仅以9.63%的体积分数,便能分别承担高达55%和42%的外加载荷。
2.4 位错密度演化
位错密度演化的分析结果见图6。在开始变形前,打印态和退火态合金中的Al基体位错密度分别为6.83 × 1014和3.85 × 1014 m-2。在塑性变形阶段,打印态合金中Al基体的位错密度变化可分为3个阶段,这与图2中划分的3个阶段相吻合。因此,可以推测位错密度的这种阶段性演化是导致打印态合金加工硬化行为呈现三阶段特征的根本原因。具体来说,阶段A的位错密度迅速上升,阶段B增长放缓,而到了阶段C,位错密度再次以较快的速度增加。对于退火态合金,随着宏观应变的增加,Al基体的位错密度最初增长较快,但随后增速逐渐放缓。在整个加载变形过程中,位错密度的变化相对平稳。在均匀变形阶段,打印态和退火态合金中Al基体的位错密度最大值分别为6.92 × 1015和1.63 × 1015 m-2。
图6
图6
打印态和退火态合金中Al基体的位错密度演化
Fig.6
Dislocation density evolutions of Al matrix of as-built and annealed samples
2.5 去应力退火对合金变形行为的影响
2.5.1 退火前后Al和Si相对合金加工硬化率的贡献
式中,
图7
图7
Al和Si两相对合金加工硬化率的贡献(这里重新绘制了图2b中的合金数据用于对比)
Fig.7
Contributions of Al and Si phases to the strain hardening rate of the alloys (The data of alloys in Fig.2b was plotted here again for comparison)
(a) as-built sample (b) annealed sample
2.5.2 退火热处理对位错行为的影响
图6表明去应力退火处理显著降低了合金的初始位错密度,这个结果与Chen等[10]的电子背散射衍射(EBSD)观察结果一致:增材制造AlSi10Mg合金在280℃退火2 h后,核平均取向差(KAM)减小,证实了位错密度的降低。同时,变形过程中的位错密度演化表明,退火态合金的位错密度增长率远低于打印态合金,这与前期研究[25]中位错密度的演化趋势相符:前期对增材制造AlSi3.5Mg2.5合金的研究表明,直接时效态和退火态AlSi3.5Mg2.5合金的初始位错密度分别为6.00 × 1014和1.40 × 1014 m-2;变形5.7%后,位错密度分别增加到2.47 × 1015和3.22 × 1014 m-2。
3 结论
(1) LPBF制备的AlSi10Mg打印态合金的加工硬化率可分为3个阶段:阶段A的加工硬化率非常高,但随应变增加而快速降低;阶段B的加工硬化率变化很小;阶段C的加工硬化率随着应变增大而逐步减小。LPBF制备的AlSi10Mg合金在270℃退火2 h后,强度和加工硬化率均降低,且加工硬化率随应变呈光滑连续降低。微观结构方面,退火态合金中Al-Si共晶网络局部断裂,Si相发生球化,且Al基体发生了回复,位错密度降低。
(2) AlSi10Mg打印态合金的Al、Si 2相的晶格应变演化可细分为5个阶段:纯弹性变形阶段、弹塑性转变前期、弹塑性转变后期、宏观塑性变形阶段和损伤阶段。对于AlSi10Mg退火态合金,Al、Si两相在加载阶段的晶格应变演化可细分为4个阶段:纯弹性变形阶段、弹塑性转变前期、弹塑性转变后期和宏观塑性变形阶段。
(3) Al{111}晶格应变的演化行为与宏观变形行为最接近,且其残余晶格应变最小,因此用Al{111}晶格应变/应力代表Al基体的整体行为比Al{311}更合适。均匀变形阶段,打印态和退火态合金中Si{311}晶格应力最大值分别达到了2885和1224 MPa,表明Si相产生了庞载荷传递强化效应。对加工硬化率的分析表明,无论是打印态合金还是退火态合金,在塑性变形前期,Si相对加工硬化率的贡献均占主导。
(4) 打印态合金的位错密度演化分为3个阶段:阶段A的位错密度迅速上升;阶段B增长放缓;而到了阶段C,位错密度再次以较快的速度增加。对于退火态合金,Al-Si共晶网络出现局部断裂,Si相发生部分球化,对位错的阻碍作用降低,随着宏观应变的增加,Al基体的位错密度最初增长较快,但随后增速逐渐放缓。
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