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Vol. 8. Issue 3.
Pages 2649-2661 (May - June 2019)
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Vol. 8. Issue 3.
Pages 2649-2661 (May - June 2019)
Original Article
DOI: 10.1016/j.jmrt.2018.12.025
Open Access
Effects of melting-mixing ratio on the interfacial microstructure and tensile properties of austenitic–ferritic stainless steel joints
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Xu Zhanga, Yu Zhanga,b, Yudong Wua, Sansan Aoa, Zhen Luoa,
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lz_tju@163.com

Corresponding author.
a School of Materials Science and Engineering, Tianjin University, Tianjin 300072, China
b Department of Mechanical Engineering, Tsinghua University, Beijing 100084, China
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Tables (5)
Table 1. Chemical composition of base metals (wt.%).
Table 2. Mechanical properties of base metals.
Table 3. Welding parameters.
Table 4. Chemical composition of the fusion zone under different offsets (wt.%).
Table 5. Strain hardening exponent calculated from the DIC test.
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Abstract

In this study, the microstructure and mechanical properties of dissimilar austenitic–ferritic stainless steel joints welded by gas tungsten arc welding were investigated. Different offsets of the tungsten electrode's tip from the centerline of the assembled workpieces were used to obtain different elemental melting-mixing ratios. In tensile tests, two failure modes were found: the weld/workpiece interface failure mode (low strength and ductility) and the AISI 430 base metal failure mode (high strength and ductility). With a large offset (≥0.5mm) toward the AISI 304 side, additional elemental Ni was melted into the welding pool, which changed the solidification mode from ferritic to ferritic–austenitic. This leads to a dramatic microstructure and local property change at the weld/workpiece interface. The local stress–strain curves of various heat-affected zones from different joints were obtained by the digital image correlation technique. This technique showed that the steep local microstructural and mechanical property transition of the dissimilar austenitic–ferritic joints will lead to failure at the weld/workpiece interface.

Keywords:
Austenitic–ferritic stainless steel joints
Digital image correlation
Local stress–strain curves
Tensile properties
Weld failure mode
Full Text
1Introduction

Welding joints between austenitic and ferritic stainless steels are extensively utilized in the chemical, petrochemical, and nuclear industries and in vacuum and cryogenic equipment [1,2]. This kind of hybrid material structure can take advantage of both the high corrosion resistance of austenitic stainless steels and the low cost of ferritic stainless steels.

Many studies have focused on the mechanical properties of dissimilar austenitic–ferritic stainless steel joints. In previous research, dissimilar austenitic–ferritic stainless steel joints have been welded by electron beam welding, laser beam welding, gas tungsten arc welding (GTAW), friction stir welding, and resistance spot welding [3–5]. Among these studies, the microstructures of dissimilar austenitic–ferritic stainless steel joints welded by different methods all present a duplex phase feature (mixing state of austenite and ferrite). It could be observed that, although the phase transition patterns are similar to each other among austenitic–ferritic stainless steel joints (processed by different welding methods), different element proportions in each local zone will lead to dramatic differences in the strength, ductility, and failure mode of the joints. Solution of this important issue will provide a comprehensive understanding of the mechanical behavior of a heterogeneous structure (and particularly a dissimilar welding joint).

Few studies have examined the local constitutive behavior of dissimilar austenitic–ferritic stainless steel joints, which are typical heterogeneous structures. Typical mechanical tests do not readily reveal the fracturing process or local properties of a dissimilar metal joint. It is necessary to develop new methods to understand this kind of complex load-bearing structure.

Digital image correlation (DIC) is a non-contact full-field method to measure the deformation contours in any material [6]. Currently, DIC is a widely used experimental technique to assess the mechanical behavior of various materials, particularly the fracture characteristics and failure modes of a variety of structural integrity components.

The DIC technique can directly and quantitatively determine the deformation field and characterize the deformation mechanism of various materials subjected to mechanical, thermal, or other loading [7–9]. Further, various mechanical parameters of a material, including Young's modulus and Poisson's ratio, can be further identified based on the computed displacement fields or strain fields.

Digital image correlation is often applied to obtain an in situ observation of welding joints. Genevois [10] and Lockwood et al. [11] applied this technique to determine the constitutive behavior of AA2024 friction stir welds. Corr et al. [12] used DIC to analyze interfacial debonding properties and fracture behavior in concrete. Leitao et al. [13] obtained local and global stress–strain curves of aluminum friction stir welds using DIC. Genevois et al. [14] also researched local and global mechanical properties of 2024 T351, 2024 T6, and 5251 O friction stir welds. Kang et al. [15] studied the characterization of constitutive behavior of dissimilar aluminum alloy resistance spot welds by implementing DIC technology.

This paper aims to characterize the microstructure and mechanical properties of dissimilar AISI 304/AISI 430 GTAW joints (austenitic–ferritic joints). The solidification paths of the austenitic–ferritic joints under different melting-mixing ratios were calculated. In addition, DIC tests were introduced to test the uniaxial tensile strength of the joints. Local stress–strain curves were obtained by the combination of force analysis and DIC. This helped provide a better understanding of the different fracturing modes triggered by different welding parameters.

2Experimental procedures2.1Welding procedures

In the present study, 1.5-mm-thick sheets of AISI 304 (austenitic stainless steel, ASS) and AISI 430 (ferritic stainless steel, FSS) were welded by GTAW under different offsets. The offset is the distance between the tip of the tungsten electrode and the boundary of the two workpieces, as shown in Fig. 1. Both sheets were 150mm long and 60mm wide. The chemical compositions and mechanical properties of the two materials are listed in Tables 1 and 2. Because the base metals have different electromagnetic properties, when the offset is 0, the arc deflects toward AISI 430, which leads to a weldment offset to the 430 side. Therefore, in this study, the tungsten electrode was offset to the AISI 304 side by various distances. The different offsets lead to different melting portions of the bilateral metal and trigger different local property distributions. Preliminary tests were conducted to determine the process parameters that affect weld formation. The process parameters are summarized in Table 3.

Fig. 1.

Schematic of the gas tungsten arc welding (GTAW) technique.

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Table 1.

Chemical composition of base metals (wt.%).

Composition  Si  Mn  Cr  Ni 
AISI 430  0.06  0.4  0.4  17.0  0.03  0.04  – 
AISI 304  0.06  0.32  1.38  18.4  0.28  0.4  8.17 
Table 2.

Mechanical properties of base metals.

Property  Ultimate tensile strength (MPa)  Yield strength (MPa)  % Elongation  Young's modulus (GPa) 
AISI 430  488  380  28  219 
AISI 304  600  250  58  194 
Table 3.

Welding parameters.

Welding materials  Welding current (amps)  Welding speed (mm/min)  Arc length (mm)  Offset (mm)  Shielding gas (L/min) 
AISI 430/AISI 304  80  200  2.5  0/0.2/0.5/1.5  11 
2.2Metallographic analysis

After welding, metallographic samples were cut from the joints to show the microstructure distribution from the top view. Then dissimilar ASS/FSS joints and similar FSS joint specimens were polished and etched with potassium metabisulfite solution (1g potassium metabisulfite, 10mL HCl, 100mL H2O) to distinguish between austenite and ferrite [16]. The microstructure of similar ASS joints was revealed by a saturated and mixed solution of FeCl3 and HCl diluted to a concentration of 40%. The microstructure was analyzed by optical microscopy.

To analyze the effect of the fusion zone (FZ) composition on the solidification mode of dissimilar austenitic–ferritic welds, temperature versus mole solid fraction (TfS) curves were calculated for each offset. The chemical composition of the FZ was simplified as a Cr–Fe–Ni ternary system. The commercial thermodynamic software Pandat [17] was used to calculate the TfS curves with the Scheil–Gulliver solidification model [18,19]. The composition of the weld was estimated according to Eq. (1)[20]:

where (wt.%E)weld, (wt.%E)ASS, and (wt.%E)FSS are the weight percentages of each element E (Cr, Fe, and Ni) in the weld, ASS workpiece, and FSS workpiece, respectively, and Vweld, VASS, and VFSS are the volume of the weld, melted ASS workpiece, and melted FSS workpiece, respectively.

2.3Mechanical testing and in situ observation

To test the mechanical properties, tensile test specimens were machined according to GB/T 2651-2008 (dimensions are shown in Fig. 2). In addition, micro Vickers hardness distribution tests were conducted on the front of the joints vertical to the welding direction by imposing loads of 200g for 10s.

Fig. 2.

Dimensions of the tensile testing specimen, sampling location for metallographic sample, and location of Vickers hardness testing line.

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For this work, a commercially available DIC algorithm was used. Vic-2D from correlated solutions is a two-dimensional (2D) displacement and strain analysis package using full-field DIC. In this algorithm, the 2D displacement field is represented by a uniform parametric B-spline surface function with unknown coefficients.

The tensile samples were speckled by paint and marking pen. A digital camera (3456×5184 pixels) captured the fracturing process with intervals of 0.2s during the tensile test. To avoid undesired shadows in the pictures, a 100-W spotlight was used (as shown in Fig. 3). The digital camera was synchronized with the other monitoring devices, such as the strain gauges. An isostress condition was assumed for the local stress calculation.

Fig. 3.

Typical optical image acquisition system for the 2D DIC method.

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3Results and discussion3.1Microstructures of the austenitic–ferritic stainless steel joints

The typical microstructure of an ASS/FSS joint is shown in Fig. 4. It was found that the joint could be divided into a FZ and heat-affected zone (HAZ). Note that the fusion lines have a different appearance between the ASS and FSS sides. Kou et al. pointed out that this phenomenon is due to the melting point difference of two workpieces and the liquid pool's flow pattern [21,22]. In this study, because the melting points of the two kinds of metal are very similar (1398–1454°C for AISI 304 and 1427°C for AISI 430), the macro-segregation zone is very narrow (according to Kou's theory). Thus, the effect of this phenomenon on the mechanical properties of the joint is neglected in the following discussion.

Fig. 4.

Typical microstructures in the ASS/FSS joint: (a) overall structure, (b) segregation zone adjacent to the ASS, (c) fusion zone, (d) GCHAZ adjacent to the FSS, and (e) GRHAZ adjacent to the FSS.

(2.36MB).

In the HAZ, the ASS and FSS do not mix. The HAZ adjacent to the ASS interface contains fully austenite grains [Fig. 4(b)]. The HAZ on the FSS side can be divided into two distinct metallurgical transformation zones, namely, grain coarser heat-affected zone (GCHAZ) and grain refining heat-affected zone (GRHAZ), as indicated in Fig. 4(d) and (e). The ferrite grain growth in the GCHAZ can be substantial due to the absence of elevated-temperature austenite [23]. In the GRHAZ, some transformation to austenite occurs along the ferrite grain boundaries upon cooling, which plays an important role in preventing the grain growth of ferrite. The difference between phase transformation of the GCHAZ and GRHAZ means that they have significantly different properties. In Fig. 4(b) and (d), the grain size of the HAZ in the 430 side is larger than that in the 304 side. Celik and Alsaran [24] have found that the HAZ of the austenitic steel side shows a slight grain recrystallization, while that of the ferritic stainless steel side indicates a clear recrystallization during welding.

The FZ of the dissimilar ASS/FSS joint sample experiences the mixing of the two workpieces, which results in a duplex phase feature at this site. When C from FSS and Ni from ASS migrates into this region, the local strength of the FZ will increase. This is because the high content of Ni increases the solution-strengthening effect of C [25].

Fig. 5 shows the FZ/GCHAZ interface of each ASS/FSS joint. Note that the epitaxial solidification phenomenon is obvious in the 0- and 0.2-mm offset samples. In contrast, a conjugation line exists between the FZ and GCHAZ in the 0.5- and 1.5-mm offset samples, which indicates the dramatic change at this site.

Fig. 5.

FZ/GCHAZ interface of each ASS/FSS joint: (a) offset=0, (b) offset=0.2mm, (c) offset=0.5mm, and (d) offset=1.5mm.

(2.74MB).

The microhardness distribution across the FZ and HAZ of different joints is presented in Fig. 6. In the weldment on the FSS side, the hardness of the FZ near the fusion line is noticeably higher than that of the other parts in the weldment. In addition, the hardness of the FZ for smaller weld offsets changes erratically, while it becomes more uniform in the large weld offset. These results are consistent with the research of Reddy et al. [3]. The GCHAZ and GRHAZ on the 430 side present different microhardness values. The hardness of GCHAZ is smaller than that of GRHAZ, which is related to the different phase transformations in the two regions.

Fig. 6.

Microhardness distribution across the FZ and HAZ of different joints.

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3.2Tensile property of austenitic–ferritic stainless steel joints under different offsets

Fig. 7 shows the engineering stress–strain curves of different joints. The stress–strain curves of dissimilar ASS/FSS joints can be divided into two types. When the offset was 0 or 0.2mm, the fracturing site was in the 430 base metal. The resultant joint had high strength and ductility, with an ultimate testing strength (UTS) of about 400MPa and elongation of about 22.7%. However, when the offset was 0.5mm or 1.5mm, the fracturing site was in the FZ/GCHAZ interface (as shown in Fig. 8). Consequently, both the UTS and joint elongation clearly decreased. The larger offset decreased the strength and ductility of the joint.

Fig. 7.

Engineering stress–strain curves and fracturing sites of different joints.

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Fig. 8.

Fracturing site of dissimilar ASS/FSS joint with 0.5-mm offset.

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Figs. 9 and 10 show the typical strain variation in the loading direction, ɛyy, for dissimilar ASS/FSS joints obtained from DIC tests. The local stress–strain curves of different joints in different regions were obtained as shown in Fig. 11. The local strain was extracted from DIC, while the local stress was calculated from the engineering stress. In Fig. 11, the intersection of local stress–strain curves and the line parallel to the x-axis is thought to reflect the strain values of different regions under the same stress (i.e., at the same time).

Fig. 9.

Contour of the absolute strain along the tensile axis, ɛyy, of dissimilar ASS/FSS joint with no offset: (a) engineering stress–strain curve of the joint and (b–h) tensile strains of corresponding moments in (a).

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Fig. 10.

Contour of absolute strain along the tensile axis, ɛyy, of dissimilar ASS/FSS joint with 0.5-mm offset: (a) engineering stress–strain curve of the joint and (b–h) tensile strain of corresponding moments in (a).

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Fig. 11.

Local stress–strain curves of different joints: (a) dissimilar ASS/FSS weld with no offset and (b) dissimilar ASS/FSS weld with 0.5-mm offset.

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Fig. 9 shows the ɛyy for the dissimilar ASS/FSS joint with no offset, which failed in the weak FSS base metal. As can be seen from Fig. 11(a), because of the lower Young's modulus and lower yield strength of ASS with respect to FSS, the local strain in the ASS base metal is larger than that in the FSS base metal during the elastic deformation stage. This also results in a larger local strain in the ASS base metal at the beginning of plastic deformation. As the stress increased to 420MPa, the local strains of the ASS and FSS base metals became equivalent, about 0.14. As the uniaxial tensile test continued, the FSS base metal began to neck down. At the same time, the tensile load no longer increased, and the local strain of ASS base metal also stopped increasing. The deformation concentrated in the FSS base metal until it finally fractured. During the whole process, the local strains in the FZ and HAZs on both sides were very low.

Fig. 10 shows the ɛyy for a dissimilar ASS/FSS joint with 0.5-mm offset, which failed in the FZ/GCHAZ interface. The larger Young's modulus and yield strength of FSS with respect to ASS resulted in larger local strain in the ASS base metal during elastic deformation and the beginning of plastic deformation. At 350MPa, the stress reached the highest peak that the FZ/GCHAZ interface could withstand. Therefore, the failure occurred in the FZ/GCHAZ interface. During the whole process, even after the failure occurred, the local strain of the FZ/GCHAZ interface was very small, about 0.01. Clearly, the fracture mode was brittle fracture. It can also be seen that the Young's modulus of the HAZ adjacent to the FSS interface is equal to that of the FSS base metal. In addition, Young's modulus in the FZ was between that of the FSS and ASS base metals.

For the distribution of strain field of dissimilar joints under uniaxial tension, the important factors are the elastic modulus, yield strength, tensile strength of the two base metals, and the zone that can bear the minimum true stress. Between the metals used in this study, AISI 430 has larger Young's modulus, larger yield strength, and lower tensile strength than AISI 304. As a result, in the elastic deformation stage, the FSS base metal has low strain. When the fracture occurs in the base metal, the plastic deformation stage can be divided into two parts. The intersection of the local stress–strain curves of the two base metals is the boundary point. Before that, the strain of the material with high yield strength (FSS) is lower, but the low tensile strength of FSS leads to a higher strain growth rate than that of ASS. After this point, the strain in materials with low tensile strength (FSS) is higher, and the strain growth rate becomes faster than that in ASS until necking and fracture occur subsequently. When the fracture occurs in the FZ/GCHAZ interface, according to the iso-stress condition, the true stress in every zone is same, so the fracture will occur in the zone that can bear the minimum true stress. If this zone is hard and brittle, a brittle fracture will occur. This study found that the tensile limit in the interface of the FZ and HAZ zone on the 430 side was clearly reduced due to the abrupt change of structure and local mechanical properties. Fracturing at this location is more thoroughly discussed in Section 3.3.

3.3Metallurgical and fracturing analysis

The tensile tests revealed two kinds of failure mode for the ASS/FSS joints under different offsets. When the offset was larger than 0.5mm, the fracture site was the FZ/GCHAZ interface, which implies an important microstructure and local property change in this region.

Increasing the offset led to more ASS workpiece melting into the welding pool, which triggers the transition of the solidification mode from ferritic to ferritic–austenitic. Table 4 shows the chemical composition of the FZ under different offsets, as calculated by Eq. (1). Fig. 12(a) presents the calculated solidification path of the ASS/FSS welds under different offsets. It can be observed that the additional Ni melted into the welding pool (caused by a large offset) changed the solidification mode significantly. In the figure, the break points on the TfS curves marked by black arrows refer to the generation of the austenitic γ phase. After taking the solid-state phase transformation of the GCHAZ on the AISI 430 side into consideration [23], the schematic of the 0- and 0.2-mm offset samples could be drawn as Fig. 12(b). A corresponding schematic for the 0.5- and 1.5-mm offset samples is shown in Fig. 12(c). For samples with small offset, only the ferritic δ phase precipitated during solidification, which transformed into the ferritic α phase at room temperature. This led to epitaxial grain growth at the interface. For samples with large offset, the ferritic–austenitic solidification mode led to a dramatic microstructure change at the FZ/GCHAZ interface.

Table 4.

Chemical composition of the fusion zone under different offsets (wt.%).

Composition  Fe  Cr  Ni 
Offset=78.71  17.66  3.53 
Offset=0.2mm  77.73  17.83  4.33 
Offset=0.5mm  76.42  18.08  5.40 
Offset=1.5mm  74.33  18.46  7.11 
Fig. 12.

(a) TfS curves of the dissimilar ASS/FSS welds, (b) schematic of the ferritic solidification mode process, and (c) schematic of the ferritic–austenitic solidification mode process.

(0.48MB).

The dramatic microstructure changes at the interface led to a sharp transition of local mechanical properties, which was revealed by the DIC test. The strain hardening exponent (n) is an important index that reflects the plastic properties of materials, and it can be calculated by Eq. (2):

where σ is the true stress in the uniform plastic deformation stage, K is the strength coefficient, n is the strain hardening exponent, and ɛ is the true strain in the uniform plastic deformation stage of the corresponding moment. The n values calculated from the DIC test are given in Table 5.

Table 5.

Strain hardening exponent calculated from the DIC test.

  ASS BM  FSS BM  GCHAZ  FZ 
Offset=0mm  0.200  0.140  0.168  0.266 
Offset=0.5mm  0.220  0.145  0.166  0.360 

As shown in Table 5, the difference of n between GCHAZ and FZ with large offset is bigger than that with little offset, which proves that a steep local mechanical property transition takes place within the sample with large offset.

Fig. 13 provides SEM images of the tensile fracture surface of two dissimilar joints. Fig. 13(a) shows the fracture surface of a dissimilar ASS/FSS joint with no offset. The presence of equiaxed dimples, which are produced by the uniaxial tensile test, confirms that the fracture is ductile fracture and occurred in the FSS base metal. Fig. 13(b) shows the tensile fracture surface of a dissimilar ASS/FSS joint with a 0.5-mm offset, which failed at the FZ/GCHAZ interface. Cleavage steps and a tearing ridge in the fracture area confirms that the failure mode is brittle fracture.

Fig. 13.

SEM image of tensile fracture surface of two joints: (a) dissimilar ASS/FSS weld with no offset and (b) dissimilar ASS/FSS weld with 0.5-mm offset.

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4Conclusion

Microstructure and mechanical properties of GTAW dissimilar ASS/FSS joints were investigated. Also, the mechanical behavior of dissimilar ASS/FSS joints was determined by DIC. The following conclusions can be drawn from this research:

  • (1)

    The solidification path of dissimilar AISI 304 and AISI 430 GTAW joints were significantly affected by the offset of the tungsten electrode during welding. The offset substantially changed the composition of the weld. Under a large offset, due to the melting of additional elemental Ni into the welding pool, the solidification mode of the weld became ferritic–austenitic. This led to a dramatic microstructure change at the interface between the FZ and HAZ in the AISI 430 side.

  • (2)

    The local constitutive behavior of different local zones in dissimilar austenitic–ferritic stainless steel joints were extracted with the DIC technique. The results for dissimilar ASS/FSS joints show that the difference of strain hardening exponent between GCHAZ and FZ with large offset is bigger than that with small offset, which is related to dramatic microstructure changes at the interface with large offset.

  • (3)

    The dramatic microstructure change and sharp transition of local mechanical properties at the FZ/GCHAZ interface results in a brittle fracture at the FZ/GCHAZ interface with large offset.

Conflicts of interest

The authors declare no conflicts of interest.

Acknowledgments

This work was supported by National Key R&D Program of China (No. 2018YFB1107900), the National Natural Science Foundation of China (No. 51575383) and Natural Science Foundation of Tianjin City (No. 18JCQNJC04100).

References
[1]
P. Bala Srinivasan, M.P.S. Kumar.
Characterisation of thin section dissimilar weld joint comprising austenitic and ferritic stainless steels.
Metal Sci J, 24 (2013), pp. 392-398
[2]
H. Luo, Z. Li, A.M. Mingers, et al.
Corrosion behavior of an equiatomic CoCrFeMnNi high-entropy alloy compared with 304 stainless steel in sulfuric acid solution.
Corros Sci, 134 (2018),
[3]
G.M. Reddy, T. Mohandas, A.S. Rao, V.V. Satyanarayana.
Influence of welding processes on microstructure and mechanical properties of dissimilar austenitic–ferritic stainless steel welds.
Adv Manuf Process, 20 (2005), pp. 147-173
[4]
U. Çaligülü, H. Dikbaş, M. Taşkin.
Microstructural characteristic of dissimilar welded components (AISI 430 ferritic-AISI 304 austenitic stainless steels) by CO2 laser beam welding (LBW).
Gazi Univ J Sci, 25 (2010), pp. 35-52
[5]
M.H. Bina, M. Jamali, M. Shamanian, H. Sabet.
Investigation on the resistance spot-welded austenitic/ferritic stainless steel.
Int J Adv Manuf Technol, 75 (2014), pp. 1371-1379
[6]
B. Pan, K. Qian, H. Xie, A. Asundi.
Topical review: two-dimensional digital image correlation for in-plane displacement and strain measurement: a review.
Meas Sci Technol, 20 (2009), pp. 152-154
[7]
S.M. Kleinendorst, J.P.M. Hoefnagels, M.G.D. Geers.
Mechanical shape correlation: a novel integrated digital image correlation approach.
(2018),
[8]
B.L. Boyce, P.L. Reu, C.V. Robino.
The constitutive behavior of laser welds in 304l stainless steel determined by digital image correlation.
Metall Mater Trans A, 37 (2006), pp. 2481-2492
[9]
G. Sun, F. Xu, G. Li, X. Huang, Q. Li.
Determination of mechanical properties of the weld line by combining micro-indentation with inverse modeling.
Comput Mater Sci, 85 (2014), pp. 347-362
[10]
C. Genevois, A. Deschamps, A. Denquin, B. Doisneau-Cottignies.
Quantitative investigation of precipitation and mechanical behaviour for AA2024 friction stir welds.
Acta Mater, 53 (2005), pp. 2447-2458
[11]
W.D. Lockwood, B. Tomaz, A.P. Reynolds.
Mechanical response of friction stir welded aa2024: experiment and modeling.
Mater Sci Eng A, 323 (2002), pp. 348-353
[12]
D. Corr, M. Accardi, L. Graham-Brady, S. Shah.
Digital image correlation analysis of interfacial debonding properties and fracture behavior in concrete.
Eng Fract Mech, 74 (2007), pp. 109-121
[13]
C. Leitão, I. Galvão, R.M. Leal, D.M. Rodrigues.
Determination of local constitutive properties of aluminium friction stir welds using digital image correlation.
Mater Des, 33 (2012), pp. 69-74
[14]
C. Genevois, A. Deschamps, P. Vacher.
Comparative study on local and global mechanical properties of 2024 T351, 2024 T6 and 5251 O friction stir welds.
Mater Sci Eng A, 415 (2006), pp. 162-170
[15]
J. Kang, B. Shalchi-Amirkhiz, Y. Chen, et al.
Characterization of constitutive behavior of dissimilar aluminum alloy resistance spot welds.
Weld J, (2016),
[16]
C. Celada.
Chemical banding revealed by chemical etching in a cold-rolled metastable stainless steel.
Mater Charact, 84 (2013), pp. 142-152
[17]
Pandat – a phase diagram calculation software package for multicomponent systems.
CompuTherm LLC, (2001), pp. 53719
[18]
G.H. Gulliver.
Metallic alloys (appendix).
Charles Griffin & Co., Ltd., (1922),
[19]
E. Scheil.
Bemerkungen zur Schlichtkristallbildung.
Z Met, 34 (1942), pp. 70
[20]
T. Yuan, X. Chai, Z. Luo, et al.
Predicting susceptibility of magnesium alloys to weld-edge cracking.
Acta Mater, 90 (2015), pp. 242-251
[21]
T. Soysal, S. Kou, D. Tat, et al.
Macrosegregation in dissimilar-metal fusion welding.
Acta Mater, 110 (2016), pp. 149-160
[22]
Y.K. Yang, S. Kou.
Macrosegregation in Al–Si welds made with dissimilar filler metals.
Sci Technol Weld Join, 13 (2010), pp. 318-326
[23]
M. Alizadeh-Sh, S.P.H. Marashi, M. Pouranvari.
Resistance spot welding of AISI 430 ferritic stainless steel: phase transformations and mechanical properties.
Mater Des, 56 (2014), pp. 258-263
[24]
A. Celik, A. Alsaran.
Mechanical and structural properties of similar and dissimilar steel joints.
Mater Charact, 43 (1999), pp. 311-318
[25]
W. Sha, Z. Guo.
Maraging steels.
first ed., Woodhead, (2009),
Journal of Materials Research and Technology

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