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匙孔内光纤激光致喷发蒸汽对焊接过程的影响

孔华, 赵振家, 邹江林, 王紫, 黄泽泓

孔华, 赵振家, 邹江林, 王紫, 黄泽泓. 匙孔内光纤激光致喷发蒸汽对焊接过程的影响[J]. 焊接学报, 2023, 44(5): 20-26. DOI: 10.12073/j.hjxb.20220530001
引用本文: 孔华, 赵振家, 邹江林, 王紫, 黄泽泓. 匙孔内光纤激光致喷发蒸汽对焊接过程的影响[J]. 焊接学报, 2023, 44(5): 20-26. DOI: 10.12073/j.hjxb.20220530001
KONG Hua, ZHAO Zhenjia, ZOU Jianglin, WANG Zi, HUANG Zehong. The influence of laser-induced plume in the keyhole on the welding process[J]. TRANSACTIONS OF THE CHINA WELDING INSTITUTION, 2023, 44(5): 20-26. DOI: 10.12073/j.hjxb.20220530001
Citation: KONG Hua, ZHAO Zhenjia, ZOU Jianglin, WANG Zi, HUANG Zehong. The influence of laser-induced plume in the keyhole on the welding process[J]. TRANSACTIONS OF THE CHINA WELDING INSTITUTION, 2023, 44(5): 20-26. DOI: 10.12073/j.hjxb.20220530001

匙孔内光纤激光致喷发蒸汽对焊接过程的影响

基金项目: 国家自然科学基金面上项目(51875007);北京市自然科学基金面上项目(3222004)
详细信息
    作者简介:

    孔华,硕士;主要研究方向为激光先进制造;Email: kongh@emails.bjut.edu.cn

    通讯作者:

    邹江林,博士,副研究员,博士研究生导师;Email: zoujianglin@bjut.edu.cn

  • 中图分类号: TG 456.7

The influence of laser-induced plume in the keyhole on the welding process

  • 摘要: 光纤激光深熔焊接羽辉由匙孔内激光致蒸汽喷发所致,对焊接过程存在严重的负面影响. 文中通过改变匙孔内激光致蒸汽的喷发特征,研究羽辉对光纤激光深熔焊接过程的影响规律. 结果表明,随着焊接速度的提高,沿焊接方向的匙孔口长度逐渐增大,匙孔前壁的倾斜角则逐渐减小. 该现象导致孔内激光致喷发蒸汽的特征发生变化:底部摆动羽辉的喷发方向逐渐沿焊接反方向偏离激光束,狭长形羽辉的高度则逐渐降低直至消失;羽辉对焊接熔深的负面影响也逐渐减小直至消失,但飞溅数量逐渐增多,焊缝表面成形则逐渐恶化. 进一步分析表明,匙孔前壁激光致蒸汽的喷发方向变化是底部摆动羽辉的喷发方向和狭长形羽辉高度均发生改变的主要原因;提高焊接速度可降低羽辉对焊接过程的负面影响,但匙孔前壁激光致蒸汽对匙孔后壁的冲击作用将导致孔口沿焊接方向的长度变大、飞溅增多、焊缝表面成形质量变差.
    Abstract: The plume can be divided into two parts: the fluctuating portion that emerges from the keyhole, called the lower fluctuating plume, and the portion that resembles a focused laser beam, referred to as the narrow plume. The changes in the morphology of these two plume parts and their influence on the welding process were studied. The results show that with increasing welding speed, the eruption direction of the lower fluctuating plume gradually deviates from the laser beam in the opposite direction of welding. The height of the narrow and long plume gradually decreases until it disappears. The effect of plume glow on the depth/width of the melt gradually decreases until it disappears. The forming quality of the weld surface gradually deteriorates. Increasing the welding speed reduces the negative impact of the narrow and elongated plume on the depth/width of the melt. The impact of the lower fluctuating plume on the back wall of the keyhole causes the length of the orifice along the welding direction to become larger, increasing spatter and reducing the forming quality of the weld surface.
  • 近年来,中国大力推进和实施清洁能源战略,由于环保政策实施力度加大,民用和工业“煤改气”超出预期,2017年冬季中国遭遇大面积严重“气荒”. 根据中国石油集团经济技术研究院发布的《2018年国内外油气行业发展报告》,2018年中国天然气进口量达1254亿立方米,同比增长31.7%, 对外依存度45.3%. 基于国内对天然气实际需求情况,急需大量建设天然气输送管线、大型的石油、天然气储罐等国家重点工程项目. 建造液化天然气(liquefied natural gas,LNG)大型储罐可以大大降低天然气储运成本[1],解决中国天然气季节性供需矛盾. LNG低温储罐是其中的关键设备,一般由外罐、保冷层和内罐组成. 06Ni9DR钢具有焊接性能良好、强度高、−196 ℃低温韧性优异等特点[2],成为建造LNG储罐的首选材料. 然而这一材料前期主要依赖于进口. 2008年以来,鞍钢等国内钢铁企业积极进行低温用06Ni9DR钢的研发,并且最终研发成功,成果显著,冲破了国外对06Ni9DR钢的长期垄断,大力推动了国内LNG工业的发展进程. 而其焊接技术是LNG低温储罐建设核心技术之一. 06Ni9DR钢焊接过程中易出现的问题一般包括热裂纹、冷裂纹、电弧磁偏吹和低温韧性不足[3-4]. 目前,国内有关国产06Ni9DR 钢焊接热影响区(HAZ)的组织和性能的研究较少. 文中采用焊接热模拟的方法研究一次焊接热循环对06Ni9DR 钢HAZ组织和低温韧性的影响,为其焊接工艺的制定提供理论依据.

    试验用钢板是国内某著名钢厂生产的06Ni9DR钢,供货状态为淬火+回火(Quenching + tempering, QT). 奥氏体化温度790-850 ℃,冷却介质水. 回火温度540 ~ 600 ℃,冷却介质空气. 其化学成分如表1所示,测得的临界转变温度Ac1Ac3分别为615 ℃ 和702 ℃.

    表  1  06Ni9DR钢化学成分(质量分数,%)
    Table  1.  Chemical compositions of 06Ni9 steel
    CSiMnPSNi
    0.05 0.180 0.64 0.004 0.001 9.12
    AlV Cu Cr Mo Fe
    0.0250.004 0.038 0.031 0.005 余量
    下载: 导出CSV 
    | 显示表格

    采用Gleeble3500热模拟试验机再现HAZ的组织. 热模拟焊接条件为16 kJ/cm. 热模拟试样尺寸为80 mm × 10.5 mm × 10.5 mm. 热模拟参数如图1所示,四种峰值温度分别代表HAZ的粗晶区(1 350 ℃,CGHAZ)、细晶区(950 ℃,FGHAZ)、不完全淬火区(680 ℃,ICHAZ)和回火区(600 ℃,SCHAZ).

    图  1  焊接热模拟曲线
    Figure  1.  Schematic diagram of welding thermal cycles

    焊接热模拟试验后,采用标准V形缺口试样进行冲击试验,尺寸为55 mm × 10 mm × 10 mm,温度为−196 ℃. 用维氏硬度计测试硬度,扫描电镜(SEM)和背散射电子衍射(EBSD)观察显微组织及晶界,采用XRD测量残余奥氏体含量,按照五峰六线法和全谱拟合计算残余奥氏体的体积和质量分数,SEM观察冲击断口形貌.

    母材及HAZ的−196 ℃冲击吸收能量和硬度值如表2所示. 母材的冲击吸收能量为134 J,硬度值为247.3 HV10. HAZ的低温冲击韧性明显降低,说明HAZ整体发生了脆化. 峰值温度1 350 ℃粗晶区的冲击吸收能量最低,为母材的15.67%,21 J. 峰值温度950 ℃细晶区的冲击吸收能量最高,为108.3 J. 峰值温度680 ℃不完全淬火区的冲击吸收能量为105 J,较细晶区稍有降低. 峰值温度600 ℃的回火区冲击吸收能量为68.3 J,为母材的50.97%. 回火区的硬度为245.42 HV10,与母材相近. 除此之外,其余HAZ硬度相差不大,在340 HV10 ~ 350 HV10范围内,但与母材相比,提高了约100 HV10.

    表  2  母材及HAZ性能
    Table  2.  Properties of base metal and HAZ
    峰值温度Tp/℃冲击吸收能量AKV(−196℃)/J硬度H(HV10)
    母材134.0247.3
    135021.0340.0
    950108.3350.06
    680105.0342.56
    60068.3245.42
    下载: 导出CSV 
    | 显示表格

    母材和热模拟HAZ的组织如图2所示. 由图2a可见,母材06Ni9DR 钢的组织为回火马氏体和少量逆转奥氏体. 板条状马氏体分布在多边形原奥氏体晶内,逆转奥氏体主要分布在晶界上.

    图  2  扫描电镜组织形貌
    Figure  2.  SEM micrographs of base metal and simulated HAZ. (a) base metal; (b) 1350 ℃; (c) 950 ℃ ; (d) 680 ℃; (e) 600 ℃ (low power); (f) 600 ℃(high power)

    粗晶区加热温度远高于Ac3,金属属于严重过热状态,奥氏体晶粒迅速长大,因此,冷却过程中晶粒粗大的奥氏体转变为晶粒粗大的板条马氏体,还有少量的残余奥氏体(图2b). 细晶区加热温度略高于Ac3,奥氏体晶粒细小,在冷却过程中,晶粒细小的奥氏体转变为晶粒细小的马氏体以及少量的残余奥氏体(图2c). 不完全淬火区加热温度处于Ac1和Ac3之间,只有部分马氏体转变为奥氏体. 由于加热时间短,加热温度不高,Ni等化学元素不能快速扩散,导致形成的奥氏体不能快速长大. 而少量未转变的马氏体继续长大,因此室温组织为晶粒尺寸不均匀的马氏体和残余奥氏体(图2d). 由图2e可见,回火区加热温度在MsAc1之间,由马氏体直接切变生成逆转奥氏体,是非扩散型转变产物. 由于组织中C,Ni,Mn等稳定奥氏体元素聚集量较高,热稳定性很高,常温中能够稳定存在,主要分布在原奥氏体晶界、马氏体束界[5]、马氏体板条间. 高倍观察(图2f),大块逆转奥氏体在冷却过程中又形成二次板条马氏体. 回火区的组织为回火马氏体和逆转奥氏体.

    图3为HAZ含有晶界的EBSD欧拉取向图和晶界角度图,蓝色为3° ~ 15°小角度晶界,绿色为15° ~ 45°大角度晶界,黄色为大于 45°有效大角度晶界. 可见,粗晶区的原奥氏体晶界主要为15° ~ 45°的大角度晶界,晶界内部不同取向板条束之间为黄色有效大角度晶界,取向差较小的板条之间为小角度晶界(图3b). 细晶区的晶粒呈多边形块状,且细化明显(图3c). 未完全淬火区晶粒大小较不均匀(图3d). 回火区的原奥氏体晶粒较母材无明显变化(图3e). 如图3f所示,45°以上有效大角度晶界含量分别为母材21.7%,粗晶区14.4%,细晶区24.4%,不完全淬火区33.5%,回火区23.1%.母材及HAZ马氏体相晶粒内部局域取向差分布如图4所示. 母材的局域取向差峰值在0.45°,粗晶区的局域取向差峰值为1.55°,细晶区和回火区的均为0.55°,不完全淬火区的取向差峰值为0.35°. 说明母材及HAZ晶粒内部存在着不均匀的位错密度堆积,除粗晶区以外三个区的位错密度低、产生的位错滑移运动较少,塑性较好[6].

    图  3  EBSD欧拉图和晶界角度分布
    Figure  3.  EBSD Euler graphs and local misorientation distribution.(a) base metal; (b) 1350 ℃; (c) 950 ℃; (d) 680 ℃; (e) 600 ℃; (f) grain boundary angle distribution
    图  4  局域取向差分布
    Figure  4.  Local misorientation distribution

    母材及HAZ奥氏体含量如表3所示. 由于奥氏体与马氏体密度不同,计算出的质量分数与体积分数在数值上有所差距,但是变化趋势相同. 相对母材,粗晶区奥氏体的含量与其相近,细晶区较之略有降低,不完全淬火区和回火区的奥氏体含量相对提高.

    表  3  母材及HAZ奥氏体含量
    Table  3.  Austenite content of base metal and HAZ
    峰值温度Tp/℃奥氏体体积分数V(%)奥氏体质量分数w(%)
    母材 1.97 0.90
    1350 1.95 0.88
    950 1.88 0.63
    680 3.10 1.28
    600 2.60 1.20
    下载: 导出CSV 
    | 显示表格

    母材断口包括纤维区、放射区和剪切唇. 放射区有细小的河流花样和撕裂棱,为准解理断裂(图5a). 粗晶区(图5b)宏观断口平齐,几乎无塑性变形,纤维区及剪切唇基本消失,几乎全部为放射区,且该区无表征微小塑性变形的放射线花样,呈现闪光小面组成的结晶状断口形貌. 另外,可见一系列取向不同的平坦、光滑的解理小刻面. 这些小刻面尺寸大致相当于马氏体束大小. 每个小刻面上有解理台阶、河流花样、二次裂纹等特征,为解理断裂. 细晶区(图5c)和不完全淬火区(图5d)断口类似,主要由纤维区、较小的放射区与剪切唇组成. 放射区由细小且浅的韧窝、河流花样组成,为准解理断裂. 回火区(图5e)断口纤维区较小,主要由放射区与剪切唇组成. 放射区为准解理断裂与二次裂纹特征,可见撕裂棱、河流花样和解理面.

    图  5  断口形貌
    Figure  5.  Fracture morphology. (a) base metal; (b)1350 ℃; (c) 950 ℃; (d) 680 ℃; (e) 600 ℃

    表2可见,06Ni9DR 钢HAZ-196 ℃的冲击能量明显低于母材,HAZ整体发生了脆化. 粗晶区脆化最为明显,其次是回火区. HAZ脆化与组织、晶粒大小、大小角度晶界和位错密度等密切相关. HAZ的组织为马氏体或回火马氏体和少量的残余或逆转奥氏体(图2). 45°以上有效大角度晶界含量为680 ℃ > 950 ℃ > 600 ℃ > 母材 > 1350 ℃(图3f). 局域取向差峰值(位错密度)为1350 ℃ > 950 ℃、600 ℃ > 母材 > 680 ℃(图4). 一般来说,有效大角度晶界越多[7-8],位错密度越低,材料韧性越好. 残余或逆转奥氏体的数量、形态和分布直接影响HAZ的低温冲击功.

    粗晶区脆化的原因包括以下几个方面:一是原始奥氏体晶粒粗大[9-11]. 经计算粗晶区的晶粒平均直径超过100 μm,根据Hall-Petch公式,晶粒直径越大,裂纹扩展临界应力增量越小,韧性越低. 另外,有效大角度晶界一般分布于原奥氏体晶界和板条束界,晶粒粗大使得有效大角度晶界较少. 而大角度晶界能够阻碍裂纹的扩展,使裂纹发生钝化而产生分支,从而消耗能量,增大韧性[10]. 二是位错密度较大,导致位错缠结、交割,使其运动阻力增大,韧性降低. 三是粗大板条马氏体的存在. 板条马氏体的板条平直细长,平行排列方向性强,解理裂纹在马氏体的板条束中可连续无阻碍贯穿,导致韧性降低[12]. 四是残余奥氏体少且不稳定. 06Ni9DR 钢具有良好的低温韧性是因为逆转奥氏体的存在. 逆转奥氏体是马氏体回火时切变产生的,而残余奥氏体是马氏体相变结束剩余的过冷奥氏体. 前者含有C,Ni,Mn等奥氏体稳定元素,−196 ℃时稳定性较高,可提高材料的韧性. 后者储存能量较高,不稳定,相对于逆转奥氏体更容易转变,对低温韧性贡献较小[13].

    细晶区的有效大角度晶界和位错密度稍高于母材,残余奥氏体量稍低于母材,但其韧性远低于且硬度远高于母材. 分析认为原因可能有两个方面:一是,细晶区的马氏体为淬火马氏体,相对母材的回火马氏体,其过饱和度、内应力较大,导致塑性低而硬度高. 二是,残余奥氏体的韧化效果低于逆转奥氏体[13]. 细晶区的硬度和韧性高于其它HAZ,主要是细晶强化的结果.

    不完全淬火区的有效大角度晶界和残余奥氏体量高于母材,位错密度低于母材,但其韧性值和硬度值与细晶区相近,韧性低于母材,硬度高于母材. 导致其韧性低于母材的原因与细晶区类似,主要是淬火马氏体和残余奥氏体的存在. 另外,晶粒大小不均匀,尤其是粗大马氏体的存在也是导致韧性降低的重要因素. 需要指出的是不完全淬火区的残余奥氏体的含量远高于母材,但韧性远低于母材,这进一步说明残余奥氏体稳定性较差,对低温韧性的贡献较小[13].

    回火区的有效大角度晶界和取向差峰值稍高于母材,逆转奥氏体量远高于母材,但其韧性远低于母材. 这主要是因为母材中原有的逆转奥氏体经过600 ℃二次回火作用,C,Ni元素继续扩散,向晶界处的逆转奥氏体聚集,原逆转奥氏体继续长大,并不断生成新的逆转奥氏体,沿晶界密集呈链状分布,降低了原奥氏体晶界的结合力,为裂纹扩展提供了通道. 另外,晶界存在的部分大块奥氏体在回火过程中转变成二次板条马氏体(图2f),形成M-A组元,也会导致韧性降低[14].

    (1) 06Ni9DR 钢粗晶区和回火区的冲击吸收能量分别为母材的15.67%和50.97%. 细晶区和不完全脆化区的冲击吸收能量相近,分别为母材的80.82%和78.36%.

    (2) 粗晶区断口为解理断裂,脆化原因主要为原始奥氏体晶粒粗大及其导致的有效大角度晶界较少,残余奥氏体量少且不稳定,以及较大的位错密度和粗大马氏体的存在.

    (3) 回火区断口为准解理断裂,其脆化的主要原因是晶界呈链状分布的大块逆转奥氏体和M-A组元的存在.

    (4) 细晶区和不完全淬火区的断口均为准解理断裂. 淬火马氏体的存在和残余奥氏体的低温稳定性差导致两区韧性低于母材. 另外,晶粒大小不均匀也会导致不完全淬火区韧性降低.

  • 图  1   羽辉及匙孔观测实验示意图

    Figure  1.   Schematic diagram of plume and keyhole observation experiment

    图  2   不同焊接速度时羽辉的形貌

    Figure  2.   The shape of the plume at the different welding speeds. (a) 2 m/min; (b) 4 m/min; (c) 6 m/min; (d) 8 m/min; (e) 10 m/min; (f) 12 m/min

    图  3   不同焊接速度下的羽辉的高度及飞溅数量

    Figure  3.   Plume height and spatters number at different welding speeds

    图  4   不同焊接速度时典型的熔池和匙孔口形态

    Figure  4.   Morphology of molten pool and keyhole under the action of illuminated laser. (a) 2 m/min; (b) 4 m/min; (c) 6 m/min; (d) 8 m/min; (e) 10 m/min; (f) 12 m/min

    图  5   匙孔形貌及熔池宽度随焊接速度的变化

    Figure  5.   Variation of keyhole morphology and molten pool width with welding speed

    图  6   不同扫描速度下的焊缝表面形貌

    Figure  6.   Surface morphology of welds at different welding speeds. (a) 1 m/min; (b) 2 m/min; (c) 4 m/min; (d) 8 m/min

    图  7   横向气帘对焊缝熔深熔宽的影响规律

    Figure  7.   The influence of transverse air curtain on the penetration and with of the weld

    图  8   小孔前壁倾斜角随焊接速度的变化规律

    Figure  8.   Variation of the tilt angle of the front keyhole wall with the welding speed

    图  9   不同焊接速度下羽辉产生机理

    Figure  9.   The generation mechanism of plume at different welding speeds. (a) High welding speed; (b) Low welding speed

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  • 收稿日期:  2022-05-29
  • 网络出版日期:  2023-03-29
  • 刊出日期:  2023-05-24

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