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1
Microstructural evolution and mechanical properties of low carbon
bainitic weld metals with various nickel contents
Gaojun Mao1,2, Rui Cao1,2,*, Jianhong Chen1,2,*, Xili Guo3 and Yong Jiang3
1 Department of Materials Science and Engineering, Lanzhou University of Technology, Lanzhou, 730050, China.
2 State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Lanzhou University of Technology, Lanzhou,
730050, China.
3 Atlantic China Welding Concumables, Inc., Zigong 643000, China.
Email: zchen@lut.cn and caorui@lut.cn
Keywords: Ni content, microstructure, strength, toughness, micro-hardness
1. Introduction
There have been ongoing developments in high-strength steel weld metal (WM) to increase strength while
maintain acceptable toughness since 1960s [1]. The mechanical properties of welds are determined by the
microstructure developed during welding process [2]. The chemical composition, austenite grain size and
cooling rate are the main factors that determine the microstructure of a weld [3]. In order to increase the
mechanical properties of low-carbon bainite welds, the selection of an appropriate flux composition plays quite
an important role to obtain a fine lath bainite (LB) and acicular ferrite (AF), which can improve the properties of
welds [4].
It is crucial to add suitable alloy elements in WM to obtain the desired properties with appropriate
microstructures. The addition of Mo to Nb microalloyed steels promotes the formation of a low angle
misorientation substructure, which at the same time enhances the strengthening effect arising from grain size
reduction and dislocation density [5-7]. Mn and Cr have a large effect on the hardenability that all promote
bainite formation at the expense of AF. However, the combined effect of these elements together may also lead
to the formation of low carbon martensite [8]. Among the alloying elements of low carbon martensite/bainite
steels, an increase in Ni content was found to be an effective way to improve both the strength and the fracture
toughness [9]. As one of vital causes, grain refinement, nevertheless, can increase both strength and toughness
of WM. Since alloying elements in WM play vital roles in microstructure evolution, including grain refinement,
an appropriate alloying strategy becomes critical to promote a desirable microstructural distribution and to
achieve specification requirements [10]. The decrease of the cooling temperature and the increment of alloying
content can reduce the mean unit size. In addition, small columnar grains were associated with a Nieq between
3.4 and 6.2% [11], however, few researches report the effect on grain size of Ni addition in WM. The tensile
properties, and more specifically the yield strength (YS), are controlled by combining different strengthening
mechanisms such as microstructural refinement, solid solution hardening, precipitation strengthening, and
dislocation hardening related to the modification of the final microstructures, from the conventional polygonal
ferrite to non-polygonal phases or bainitic structures [12]. Some authors have reported an effect of the presence
of coarse grains on the ultra-tensile strength (UTS), but their effect on the YS is reported to be more limited [13].
These studies have emphasized the effects of some alloying elements, rather less attention has been paid to
effects of Ni element on properties of low carbon bainite WMs. The objective of this work is to detail the effect
of Ni content in metal powder-cored wire on the microstructure and mechanical properties of WMs obtained by
the multi-pass welding process.
2. Experiments
The Q345 HSLA steels were automatic Gas Metal Arc welded using a single consumable with various Ni
contents in the metal power flux-cored wires. Three WM specimens with different levels of Ni in WMs were
made under the same welding conditions. All of the mentioned assignments were carried out in the Atlantic
China Welding Concumables, Inc, Zigong. The chemical compositions of WMs are given in Tables 1. The
physical conditions for the welding process are listed in Table 2. The schematic illustration of the Y-type joint
by multi-pass welding is shown in Fig. 1(a), and cross-sectional views of multi-pass welding shows that the
former pass is the back sealing welding in Fig. 1(b). As the deposited WM specimens were obtained by multi-
pass welding, the microstructure in each pass differs. In this work the microstructure were divided into two
typical zones i.e. the as-deposited zone (Zone I in Fig. 1(b)) and reheated WM zone (Zone II in Fig. 1(b)),
respectively.
2
Table 1 Chemical compositions of the WM (wt. %)
Weld Specimen
C
Mn+Si
S+P
Cr+Mo
Ni
V+Cu
Ni0
0.045
1.92
0.018
1.56
0.01
0.019
Ni2
0.053
1.9
0.02
1.58
1.96
0.019
Ni4
0.052
1.87
0.021
1.55
3.65
0.18
Table 2 Welding parameters of the welding processes
Voltage
(V)
Current
(A)
Deposition rate
v/kg·h-1
Electrode feed
rate (mm/s)
Welding
speed (mm/s)
Heat input
E /kJ·mm-1
29
240
22~68
25.2
26.2
2.0 (η=0.9)
Fig.1 (a) Schematic illustration of the Y-type joint by multi-pass welding, (b) Cross-sectional views of multi-pass welding
After welding, metallographic specimens were etched by a solution of 3 pct nitric acid in ethanol to
observe detailed microstructures using SEM (FEI quanta 450). Tensile specimens were cut from the WMs along
the weld longitude direction shown in Fig. 1(a). Dimensions of round tensile specimens are shown in Fig. 2(a).
Tensile tests were carried out by the tester SHIMADZU AG-10T at room temperature (RT), -50oC, -80oC, -
110oC and -196oC. Charpy V specimens were cut from the WM that is vertical to the weld direction shown in
Fig. 1(a). Charpy V impact tests were carried out using a 450-J instrumented pendulum impact tester for
specimens at RT, -50oC, -80oC, -110oC and -196oC. Standard Charpy V specimens are used, as shown in Fig.
2(b). Microhardness was measured using a diamond indentor (HVT-1000A) with a 0.5N load lasting for 15s. On
every surface, the hardness along 10 lines with 50 points was measured. The distance between each two lines
were set to 0.5mm to make the distance of two measured point larger than 0.5mm. Additionally, measurements
taken at progressively further spacings than 0.5mm showed no appreciable change in the microhardness values,
indicating a negligible interaction between the indents.
Fig. 2 (a) Dimensions of round tensile specimens, (b) Dimensions of standard Charpy V specimens, (c) Schematic of paths measuring
micro-hardness.
3. Results and Discussion
3.1 Analysis of microstructure in WM
The microstructures of final pass as-deposited zone (also unreheated zone) and the microstructures of
reheated zone are shown in Fig. 3(a-c) and Fig. 3(d-e), respectively. Much more attention were paid to
investigating the unreheated zones in WMs, due to the fewer regions reheated by the following weld pass and
microstructures in two zones varying little in this experiment. The predominant microstructure of all specimens
is bainite, though a mixture of bainite and ferrite was observed in the SEM images. Specific analyses are
followed: Granular ferrite (GF) consists of sheaves of elongated ferrite with low misorientations and a high
dislocation density, sometimes containing roughly equiaxed islands of M/A micro-constituents (Fig.3 (a) and
(d)). It is frequently very difficult to distinguish QF grains from often ragged GF sheaves using OM and even
TEM, for the reason that both QF and GF contain dislocation substructure features [14]. In our study, both QF
and GF are considered as granular bainite (GB), which is in accordance with the viewpoint in Ref. [15]. When
M-A constituents lying on the ferrite are distributed in parallel, they turn to interval linearly arranged
intermittent short strips as shown in Fig. 3(b) and 3(e), it would be called degenerate upper bainite (DUB),
which is different from the traditional bainite. DUB reduces the existence of microcrack on condition that
carbide has not constantly precipitated along austenite grain boundaries or between bainite laths, which makes
its strength match toughness best [16]. AF nucleates at indigenous non-metallic inclusions and grows in prior
austenite grains competing with the growth of Widmanstatten ferrite and bainite nucleating at grain boundaries
[17]. And AF laths/plates partitioned large prior austenite grains into many finer and separate regions, which
(b)
(a)
(a)
(b)
(c)
3
consists of fine grained mixed microstructure with a small proportion of AF grains embedded in fine bainite
plate packets (Fig. 3 (c)) [18-20]. As for LB shown in Fig. 3(c) and 3(f), it is made up of LB ferrite ranging side
by side, and RA (retained austenite) distributing between the laths with the shape of thin slice, which is very
beneficial to improve the toughness of material [21], and the cleavage fracture resistance of bainitic
microstructures is closely related to both prior austenite and bainite packets [22]. In short, the microstructure is
changed from granular bainite (GB), degenerate upper bainite (DUB) to lath bainite (LB) and retained austenite
(RA) with the increase of Ni content from 0, 2% to 4%.
Fig.3 Microstructure in WMs—(a)Ni0, (b)Ni2, and (c)Ni4
3.2 Analysis of prior austenite grain size in WM
By using laser confocal scanning microscopy (LSCM), the specimens polished without etching were heated
to peak temperature 1350oC and then cooled at a cooling rate of 8oC/s, the relieves were produced and exhibited
in Figs. 4(a-c), where original austenite grains on the process of stimulating weld heat-cycle were observed.
These austenite grains kept original sizes in WMs basically. All the grains assumed to be spherical in shape, 8
regions similar to Figs. 4(a-c) were measured to obtain the final data in Fig. 4(d) and the sizes of grains
presented normal distribution. In conclusion, the prior austenite grain refinement can be found owning to Ni
addition. The bainite morphology depends apparently upon the austenite grain size before the transformation.
The growth of bainite sheaves is restrained by austenite grain boundaries [23]. The lath-like or plate-like AF
grains can divide large austenite grain into smaller separate regions [24]. The bainite transformed at lower
temperatures is thus confined in the smaller regions, resulting in the formation of fine-grained mixed
microstructures [20]. Accordingly, the mean and standard deviation equivalent diameters of grains in Ni4
specimen were smaller than that in any other specimen, as shown in Fig. 4(d), which means the grains in Ni4
specimen are the smallest.
Fig.4 In-situ observation of prior austenite grains of (a)Ni0, (b)Ni2 and (c)Ni4, (d) Statistics of variation trend of average grain size
3.3 Analysis of mechanical properties in WM
Tensile test results at various temperatures are depicted in Figs. 5(a) and 5(b). Every datum was averaged
from that of three specimens. YS and UTS rise with the increase of Ni content and the decrease of test
temperature. As depicted in Fig. 5(a) and 5(b), Ni0 specimen shows lower YS and UTS by 50～80MPa than that
of Ni2 specimen, while the strengths of Ni4 specimen is significantly increased by 70～120 MPa compared with
Ni2 specimen. The aforementioned results give rise to support the view in Ref. [25] that Ni gives rise to strength
due to grain refinement strengthening. Generally, the smaller the grain size, the stronger the material. In the case
of the Hall-Petch effect, decreasing the grain size can impede dislocation movement and play as one role to
increase YS [26].
4
The hardness tests were carried out along the path displayed in Fig. 2(c). Along path (1) where the hardness
values were tested from bottom weld to cap weld, and 50 points in total were made, the interval of which was
also 0.5 mm. The hardness tests were also conducted along paths (2), (3) and (4), therefore the average values
were calculated to describe hardness. And the final results are shown in Fig. 5(c). As shown in Fig. 5(c), the
average micro-hardness value reaches Hv247.7 with Ni barely, the microstructure of which consists of GB plus
QF. As the fraction of Ni increases to 2%, it adds up by Hv46.4 for the formation of large amount of DUB.
Further still, it would arrive at the value of Hv339.2 with the amount of 4% Ni as a result of majority of LB
transferred from DUB.
Figure 5(d) displays the impact toughness transition curves of three specimens. The ductile-brittle
transition temperature (DBTT) is defined as the temperature where the toughness equals the half value of the
sum of the values on the upper shelf and lower shelf, and it characterizes the transition of ductile rupture to the
brittle cleavage fracture [27]. For specimens with Ni mass fraction of 0, 2% and 4%, the values of DBTT are
about -50oC, -60oC and -100oC, respectively. Thus, the toughness transition temperature reduces with the
increase of Ni content. This result keeps in accordance with the view in Ref. [28] that grain refinement is one of
effective means for lowering the ductile-brittle transition temperature of materials. Low-temperature toughness
of Ni4 specimen is better than that of two else specimens since grain refinement can also decrease microcrack
size [25], however, room temperature of Ni0 specimen is the highest, the next is Ni4 specimen and the lowest is
Ni2 specimen.
Fig.5 (a) Variation curves and measured values of σy, (b) Variation curves and measured values of σb, (c) Average microhardness measured
weld through different paths, (d) Impact toughness transition curves measured in specimens with Ni mass fraction of 0, 2% and 4%.
The fracture of all specimens at -196oC consisted of cleavages facet with microcrack shown in Figs. 6(a-c).
According to the Griffith crack propagation theory [29], the coarse hard phase assists microcrack initiation by
reducing the initiation energy, especially when they are aggregated. For the dispersed and granular phase, the
microcrack size can be roughly regarded as the maximum diameter, and the microcrack initiation energy
decreases with increase in diameter. More importantly, reduction was presented on both the dimension of
cleavage facet, which was defined the farthest distance from crack initiation site to the boundary of cleavage
facet and the parallel adjacent distance among tear ridges regarded as lath width affecting the bainite packet size
[30]. According to Ref. [27], the dimension of cleavage facet keeps linear with the maximum sizes of austenite
grain and bainite packet. Hence, Ni addition could increase low-temperature toughness through reducing
austenite grain and bainite packet sizes. Figures 6(d-f) depict the fracture of all specimens at -80oC with small
quasi-cleavage facets, a number of small dimples, and tear ridges. The dimples were absent in Ni0 specimen, the
brittle cleavage facets dominated the fracture (Fig. 6(e)). With the increasing of Ni content, the fracture surfaces
of both Ni2 and Ni4 specimens are made up of a lager square of cleavage facets, accompanied with small and
uniform dimples (Fig. 6(e) and (f)). Likewise, the decisive factors of impact toughness here can be concluded
with the decreasing of cleavage facet square and the increasing amount of dimples due to Ni addition. In Figs.
6(g-i), not distinct difference was presented among three specimens, the fracture of which is made up with
scattered dimples of uneven sizes. On the one hand, the size and depth of dimples in Ni4 specimen are wholly a
bit larger than that in Ni2 specimen, on the other hand, a great deal smaller than that in Ni0 specimen. In other
word, it is obvious that the dimples show flatly in Fig. 6(h). For another, the varying trend of inclusions
distribution density as well as keeps pace with the above depicted varying pattern. Thus, RT toughness of these
bainite WMs can be attributed to the sizes of dimples, the depths of dimples and the distribution of inclusions.
(b)
(c)
(d)
5
Fig.6 SEM fracture graphs of (a)Ni0, (b)Ni2, and (c)Ni4 specimens at -196oC, (d)Ni0, (e)Ni2, and (f)Ni4 specimens at -80oC, (g)Ni0,
(h)Ni2, and (i)Ni4 specimens at RT.
4. Conclusions
The present study focused on the effect of Ni content on the structure–mechanical property relationship.
The primary conclusions are as follows:
(1) The predominant microstructure of weld metal was changed from GB+QF, DUB to AF+LB and austenite
grain size was refined with the increase of Ni content.
(2) Ni addition results in significant increases in YS, UTS, microhardness and low-temperature toughness whilst
a drop in DBTT, and Ni addition shows little effect on room-temperature toughness.
(3) For the weld metal with 4% Ni at optimized welding regimes, the microstructure of LB+AF can be obtained,
the satisfied mechanical properties can be achieved as follows: YS >800 MPa (RT), UTS >1000 MPa (RT),
impact enengy>60J (-80oC) and ductile-brittle transition temperature nearly at -60oC.
5. Acknowledgement
This work was financially supported by National Nature Science Foundation of China (Nos. 51675255).
References
[1] Graf M.K., and Hillenbrand H.G., High-strength large-diameter pipe for long-distance high pressure gas
pipelines, International Journal of Offshore and Polar Engineering, 14, 117-121 (2004).
[2] Joarder A., Saha S.C., Ghose A.K., Study of submerged arc weld metal and heat-affected zone
microstructures of a plain carbon steel, Welding Journal Research Supplement, 70, 141-146 (1991).
[3] Lanceter J.F., Metallurgy of Welding, Alden Press Ltd., Chapter 2, (1980).
[4] Cao R., Zhang X.B., Wang Z., Peng Y., Du W.S., Tian Z.L., Chen J.H., Investigation of Microstructural
Features Determining the Toughness of 980 MPa Bainitic Weld Metal, Metallurgical and Materials Transactions
A, 45, 815-834 (2014).
[5] Go′ mez M., Rancel L., and Medina S.F., Assessment of austenite static recrystallization and grain size
evolution during multipass hot rolling of a niobium-microalloyed steel, Metals and Materials International, 15,
689-699 (2009).
[6] Isasti N., Jorge-badiola D., Taheri M.L., Lo′ pez B., and Uranga P., Effect of composition and deformation
on coarse-grained austenite transformation in Nb-Mo microalloyed steels, Metallurgical and Materials
Transactions A, 42, 3729-3742 (2011).
[7] Isasti N., Jorge-badiola D., Tatheri M.L., Lo′ pez B., and Uranga P, Phase transformation study in Nb-Mo
microalloyed steels using dilatometry and EBSD quantification, Metallurgical and Materials Transactions A, 44,
3552-3563(2013).
6
[8] Bose-Filho W., Carvalho L.M., and Stanwood M, Effect of alloying elements on the microstructure and
inclusion formation in HSLA multipass welds, Material Characterization, 58, 26-39 (2007).
[9] Park S.G., Lee K.H., Min K.D., Kim M.C. and Lee B.S., Characterization of Phase and Misorientations on
Tempered Bainitic/Martensite Ni-Cr-Mo Low Alloy RPV Steel with Various Ni Content, Metals and Materials
International, 19, 49-54 (2013).
[10] Wang W., Liu S., Alloying and microstrucural management in developing SMAW electrodes for HSLA-
100 steel, Welding Journal, 81, 132-145(2002).
[11] Zhang Z.Y. and Farrar R.A., Columnar grain development in C-Mn-Ni low-alloy weld metals and the
influence of nickel, Journal of Material Science, 30, 5581-5588(1995).
[12] Nerea I., Denis J.B., Mitra L. Taheri, and Pello U., Microstructural features controlling mechanical
properties in Nb-Mo microalloyed steels. Part I: Yield strength, Metallurgical and Materials Transactions A, 45,
4960-4966 (2014).
[13] Patra S., Hasan SK.M.D., Narasaiah N. and Chakrbarti D. Effect of bimodal distribution in ferrite grain
sizes on the tensile properties of low-carbon steels. Material Science and Engineering A, 538, 145-155 (2012).
[14] Cizek P., Wvnne B.P., Davies. C.H.J., Muddle B.C. and Hodgson P.D., Effect of composition and austenite
deformation on the transformation characteristics of low-carbon and ultralow-carbon microalloyed steels,
Metallurgical and Materials Transactions A, 33, 1331-1349 (2002).
[15] Zhu K. Y., Barbier D. and Iung T., Characterization and quantification methods of complex BCC matrix
microstructures in advanced high strength steels, Journal of Material Science, 48, 413-423 (2013).
[16] Chaudhari V., Ritzmann H.P. and Wqllmitz G.et al, German gas pipeline first to use new generation
linepipe. Oil Gas Journal, 1, 40-47 (1995).
[17] Babu S. S. and Bhadeshia H. K. D. H., Stress and the acicular ferrite transformation. Material Science and
Engineering A, 156, 1-9 (1992).
[18] Wan X.L, Wei R. and Wu K.M., Effect of acicular ferrite formation n grain refinement in the coarse-
grained region of heat affected zone, Material Characterization, 61, 726 -731 (2010).
[19] Wan X.L., Wu K.M., Huang G., Wei R. and Cheng L., In situ bservation of austenite grain growth behavior
in the simulated coarse-grained heat-affected zone of Ti-microalloyed steels, International Journal of Minerals,
Metallugy and Materials, 21, 878-893 (2014).
[20] Wan X.L., Wu K.M, Huang G. and Wei R., In situ observations of the formation of fine-grained mixed
microstructures of acicular ferrite and bainite in the simulated coarse-grained heated-affected zone. Steel
Research International, 85, 243-250 (2014).
[21] Bhadeshia H. K. D. H., Some phase transformations in steels, Materials Science and Technology, 15, 22-29
(1999).
[22] Lambert-Perlade A., Gourgues A. F. and Pineau. A., Austenite to bainite phase transformation in the heat
affected zone of a high strength low alloy steels. Acta Materialia, 52, 2337-2348 (2004).
[23] Liu F., Xu G., Zhang Y.L., Hu H.J., Zhou L.X., and Xue Z.L., In situ observations of austenite grain
growth in Fe-C-Mn-Si super bainitic steel, International Journal of Minerals, Metallurgy and Materials,
20,1060-1069 (2013).
[24] Wan X.L., Wei R., Wu K.M., Effect of acicular ferrite formation on grain refinement in the coarse-grained
region of heat-affected zone, Material Characterization, 61, 726-731 (2010).
[25] Meyers M.A. and Chawla K.K. Mechanical Behavior of Materials, Cambridge University Press, Chapter 4,
2009.
[26] Pike L.M., Chang Y.A., Effect of Ni on vacancy concentrations and hardness in Fe-Al alloys, Metallurgical
and Materials Transactions A, 29, 1911-1915 (1998).
[27] Cao R., Li J., Liu D.S., Ma J.Y., Chen J.H., Micromechanism of decrease of impact toughness in coarse-
grain heat-affected zone of HSLA steel with increasing welding heat input, Metallurgical and Materials
Transactions A, 46, 2999-3014 (2015).
[28] Sleboda T., Muszka K., Majta J., Hale P. and Wright R.N. The possibilities of mechanical property control
in fine grained structures, Journal of Materials Processing Technology, 177, 461-464 (2006).
[29] Hu J., Du L.X., Wang J.J., Effect of welding heat input on microstructures and toughness in simulated
CGHAZ of V-N high strength steel, Material Science and Engineering A, 577,161-168 (2013).
[30] Chen J.H. and Cao R., Micromechanism of Cleavage Fracture of Metals, Elsevier, Chapter 5, 2014.
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