Proceedings of Sino-Swedish Structural Materials Symposium 2007

Effect of Welding Heat Input on Simulated HAZ

Microstructure and Toughness of a V-N Microalloyed Steel

XU W.W.’, WANG Q. E2, PAN T.2, SU H.2, YANG C.E2 ( 1 .State Key Laboratory of Metastable Materials Science and Technology. Yanshan University, Qinhuangdao 066004,

China; 2.Division of Structural Material Research, Central Iron and Steel Research Institute, Beijing 10008 1 ,China)

Abstract: This study investigates the correlation among the welding heat input, microstructure and impact toughness of the

simulated c o m e grain heat-affected zone (CGHAZ) of a V-N microalloyed steel which contained 140 ppm of nitrogen

using Gleeble simulation technique, light microscope, electron microscope and Charpy-V-Notch (CVN) testing. CVN

toughness estimated at -20°C supported that a relatively low energy input is recommended for welding of low carbon V-N

microalloyed steel, according to a fact that an impact energy as high as of more than 180j attained in the CGHAZ with heat

input less than 25Kj.cm-I, while a CGHAZ with impact energy less than 80j was obtained at a heat input more than 40 Kjan-’ . Quantitative metallography of samples showed that CVN toughness decreased with the increase of the amount of

grain boundary ferrite (GBF). Quantitative metallography of two-stage electrolytically etched samples showed that the

averaged size and ratio of lath martensite-austenite (M-A) constituent increased with increasing heat input, while the linear

density and amount of M-A decreased. It is indicated that the toughness of CGHAZ was severely dependent on the

characteristic parameters of M-A and the amount of GBF. Microhardness of M-A constituents estimated by nanoindentation

technique were much higher than that of the neighboring matrix. Severer brittle cracking susceptibility can occur in the

CGHAZ with higher heat input based on Chen et al’s suggestion that the stress concentration and triaxiality of the

neighboring matrix are increased by the hard phase particles such as M-A constituents.

Key words: phase transformation; M-A constituents; toughness; welding; HAZ;V-N microalloyed steel

1. Introduction The VN microalloying technique has been

extensively employed in the constructional steels for fabrication of heavy-section oil pipeline and H-profiled bar, bridge, heavy-duty rail, automobile beam, and so on “I . The solubility of nitrogen in austenite is much higher than that of carbon. In addition, nitrides in the steel are much more stable and thus less soluble than carbides, and their solubility in femte is considerably less than that in austenite. Hence, when nitrogen is added into the vanadium microalloyed steels, nitrides often exhibit in small particles and dispersed distributeion. Precipitation of V(C,N) and/or VN in nm size can effectively increase the strength of steel in the case of reducing carbon content by which can improve weldability obviously ” I . However, the V-N microalloyed steels contain more than two of the common microalloying additions, and the mutual

solubility of the microalloying elements in carbonitrides complicates the thermodynamics and kinerics of precipitation in the steel. This problem must be emphasized on welding of this type of steel, since the rapid welding thermal cycle can lead to dissolution followed by re-percipitation of the same or different precipitate species close to the peak temperature of thermal cycle. Under these conditions, a major proportion of the nitrides may be dissociated so that they lose their pinning function 13]. Nitrogen, which is released simultaneous, forms “free” nitrogen. During the subsequent continuous cooling, the form of vanadium nitride is changed with different microalloying system (C, N, Nb, Ti etc) and energy input, making complicated effect on CGHAZ microstructure and mechanical properties. This study investigates the correlation among the welding heat input, microstructure and impact toughness of the

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simulated coarse grain heat-affected zone (CGHAZ), and offers the basis for establishment of welding process. 2. Materials and Experimental Procedures

After melting in a 25 kg vacuum melt induction furnace, the experimental steel with chemical compositions given in n b l e 1, were hot forged into the rods of 2Omm in diameter.

Gleeble3500 thermal simulator was used to simulate single pass welding thermal cycle of CGHAZ, and the welding thermal cycle curves were made by HAZ software Package. The required parameters are heating rate at lOO"C/s, peak temperature at 1345 "C, preheating temperature at 20°C and different tg/5: 6s, lOs, 15s, 30s, 60s and 100s. Samples for optical observation were prepared by conventional grinding and polishing techniques and etched in 4% nital solution and Leper solution by which only M-A constituents can be etched. The size and amount of M-A constituents and the amount of GBF were measured by quantitative metallography using Image-Pro Plus. Further analysis on microstructure of CGHAZ was carried out using transmission electron microscopy (TEM). Microhardness of M-A

constituents and the neighboring base material were estimated by nanoindentation technique. Standard impact specimens with the size of 10 mmxlOmmx55 mm were prepared and Charpy-V-Notch impact testing was carried out at -20°C. 3. Results 3.1 Microstructures of CGHAZ

The optical microstructures obtained in different welding process were showed in Fig 1. As shown in Fig.1, in the case of tg/5=6s, the microstructure of simulated CGHAZ consisted of dominant intragranular femte (IGF) and a little amount of pearlite (P). In the case of t,=lOs, grain boundary femte (GBF) began to appear. As the tg/5 was more than 15s, microstructure of CGHAZ was composed of GBF, IGF, granular bainite (GB) and pearlite (P). Quantitative metallography result of these four micmstructures as showed in Fig.2, indicated that the amount of IGF and GB decreased significantly, while the amount and size of GBF increased with an increase in tg/5. The femte shaped in net formed along the prior austenite grain boundaries as a higher heat input was imposed. 3.2 M-A constituent in CGHAZ

The optical microstructures obtained in different

Table 1 Chemical compositions of steel used in experiment (wt%)

C Mn N V Si Mo Ti Cr B S

0. 06 1.5 0.012 0.051 0.24 0.25 0.011 0.014 0.0005 0.006

Fig. 1 Optical micrographs of CGHAZ in dflerent ta5: 6s (a), 15s (b), 30s (c) and 60s (a)

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Proceedings of SineSwedish Structural Materials Symposium 2007

Fig. 2 Area fraction of micmtructure constituents

in CGHAZ as a function of tu5

welding processes etched with Leper solution were showed in Fig.3. The white island-shaped microstructures, which only could be observed, were M-A constituents. The quantitative metallography was used to measure the parameters which described the M-A constituents in the CGHAZs with different t8/5. There exhibited two different types of M-A constituents in morphology: lathy and massive respectively. According to an early work '41, the lathy M-A constituents was defined in case of the length-to-width ratio, Ra more than 3, while the massive M-A constituents referred to Ra less than3. Then, Ra was used to characterize the morphology of M-A constituents in the present work and the ratio of lathy M-A constituents to massive ones was estimated in CGHAZ with varied t815.The averaged area of M-A constituents in pn' indicated their size, and the amount of M-A constituents in per unit length (linear density) was used to characterize their number distribution. The

6

.I 8' 4 z

0 20 10 m tn.

Fig. 3 M-A constituents in CGHAZ with dif€erent tw5: 6s (a), 15s (b), 30s (c) and 60s (d)

area fraction of M-A constituents was also estimated as a function of t8/5.

The result of quantitative metallography was showed in Fig 4. Compared with other condition of tu5, in the case of t8/5 less thanlos, there formed a large amount of massive M-A constituents in dispersed distribution and the maximum area fraction of M-A constituents was obtained with averaged size less than 2pm2 indicating very fine M-A islands presented. The average area and length-to-width ratio of M-A constituents increased with the increase of t8/5, while the number density and area fraction of M-A constituents declined.

The M-Aconstituent in CGHAZ with t8,5=30s was further observed by bright and dark field of TEM, as shown in Fig 5. It was indicated that the massive M-A

I I n a (D m

c..s

Fig. 1 Morphology, size (a), number density and area fraction (b) of M-A constituents in CGHAZ as a function of tus

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Proceedings of Sino-Swedish Structural Materials Symposium 2007

constituents with size in 2 - 3 p and highly dislocated twin structures exhibited in the grain boundaries of prior austenite.

Nanoindentation technique was used to examine the microhardness of M-A constituents and the neighboring base material in the CGHAZ with different t8/5 of 60s and 100s. The averaged value of microhardness was obtained by estimating five M-A islands for each sample. The results were shown in Table 2. Since the size of M-A constituents in CGHAZ with t8/5<30 was generally less than 2-3pm, most of the indentations exceeded the size range of M-A islands. Unfortunately, the reliable measurement of microhardness in M-A constituents failed. Anyway,

from the Table 2, in case of t8/5 more than 60s, the microhardness of M-A constituents was 1.5-2 times than that of neighboring base material.

Fig. 5 TEM observations of M-A constituents: bright field

(a) and dark field (b)

Table 2 Microhardness of M-A constituents and neighboring base matend ( m a )

M-A constituents base material

60

100

505.49

508.81

296. 11

331.63

3.3 impact toughness in the simulated HAZ The results of CVN impact toughness

measurement of the simulated CGHAZ with different t8/s performed at-20 "C were given in Fig. 6. A maximum CVN toughness as more thanl80J was obtained in case of t8/5 less than 10s and moderate toughness ranged from 60 tol20J for t8/5= 15-30s. However, when tg15 increased to 60s and above, the CVN impact toughness of CGHAZ was deteriorated by decreasing to about 205. Thus, a relatively low energy input, as less than 20kj/cm for example, would be favorable for the welding of V-N microalloyed steels. This recommendation is associated with the microstructural change in the CGHAZ with different t815, and the further discussion was given in the following. 4. Discussion 4.1 Effect of GBF on toughness

The CVN impact toughness varied with area fraction of grain boundary femte was shown in Fig 7. It was shown that the CVN impact toughness of CGHAZ with different tSt5 decreased significantly with the

- 0 20 10 60

t@, Fig. 6 Relation of CVN impact toughness of CGHAZ

at -20 "C to ks

increase of GBF. Thus, a high area fraction of GBF severely weakened the toughness of CGHAZ. This conclusion was confirmed by other work "I.

Actually, the existence of grain boundary fenite declined the impact toughness of granular bainite apparently in two different ways. On the one hand, since the strength of massive grain boundary ferrite is relatively low in comparison to that of other microstructrual constituents as pearlite, granular

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Proceedings of Sino-Swedish Structural Materials Symposium 2007

wc ' I

h a ratio of GBF. 11

Fig. 7 CVN impact toughness as a fiinction of anxi fraction

of grain boundary ferrite in CGHAZ

bainite and IGF, it's easy for the local internal stress to reach its strength limit which results in microcracking initiated and propagated inside the grain boundary femte. When a crack extended to granular bainite, it would not prevent the crack from growing strongly because the strength of matrix in granular bainite is also low. Based on the above discussion, we suggest that ong of the most important reasons for declination of toughness is that the cracking is apt to initiate and propagate in grain boundary femte. On the other hand, the declination of toughness is related to the brittleness of martensite which is transformed from austenite islands rich in carbon. When GBF starts to create in CGHAZ on early stage of 7 4 a reaction, the austenite islands rich in carbon also form simultaneously. Transformation of a high volume of GBF is generally accompanied by austenite islands with high carbon content which lead to martensite with high brittleness in M-A constituents. This is the other reason GBF declines the toughness of granular bainite. 4.2 Effect of M-A constituent on toughness

As reported by other work [61, the toughness of granular bainite was generally dependent on bainitic femte and structures in islands as martensite, M-A constituents, carbides or their combination and therefore. the morphology, fraction, number density and size of M-A constituents made great effect on the toughness of CGHAZ. Most of researches have shown

that the decreased size and number density of M-A constituents, ratio of lathy M-A constituents and area or volume fraction of M-A phase, was corresponding to the improved impact toughness of CGHAZ. However, Fig. 4(a), Fig. 4(b) and Fig. 6 indicated that the linear number density and area fraction of M-A constituents decreased with the increase of t8/5 which was accompanied by the decreasing impact toughness. This inclination was in disagreement with some early works, although the size and length-to-width ratio of M-A constituent increased with the increase of t8 /5 which was accordance with the decreased impact toughness, In summary, we suggest that these parameters used to describe the morphology of M-A constituents cannot be solely applied to understand substantially the effect of M-A constituent on toughness.

S. Aihara etc "' suggested that the mechanism of M-A constituents in influencing the toughness is that the microcracking inclines to initiate in the interface between M-A constituent and the adjacent matrix and propagate toward the base material. An earlier work by Chen et al. "I proposed that the stress concentration and triaxiality of the neighboring matrix were increased by the hard-phase particles such as the M-A constituents and that the stress in the neighboring matrix could be elevated to several times the average stress. This made it difficult for the matrix to deform plastically and the elevated stress could not be relieved effectively. Thus a concentrated deformation might occur in the interface between the matrix and the neighboring M-A constituents and the extremely high stress concentration and triaxiality could not be relieved. On this account, a brittle fracture could occur from the M-A constituents along the relatively weak boundaries and propagate to the matrix.

According to the above discussion, we can conclude that the difference of hardness between M-A constituents and the neighboring base material is an important criterion of brittle-fracture trend, In this work, we find that with the increasing tg/5 of the welding.therma1 cycle, the area fraction and number density of austenite islands rich in carbon greatly decreased as a result of a large amount of transformed

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GBF on the early stage of y-*a reaction though, a much higher hardness of M-A constituents could be expected after they further transformed. This indicated that a brittle-fracture inclination of CGHAZ was increased. 5. Conclusions

When welding of V-N microalloyed steel was carried out, the amount of intragranular femte and granular bainite in the CGHAZ decreased notability with the increasing tw5, while the amount and size of grain boundary ferrite exhibited an opposite inclination. Transformation of a large amount GBF in the CGHAZ resulted in the decreased number density and area fraction of M-A constituents, but the M-A constituents with much higher microhardness was obtained as a result of high heat input. A considerable difference in microhardness between the M-A constituents and the neighboring matrix increased the inclination of brittle fracture occurred in GCHAZ. With all the above results considered, a relatively low heat input is recommended

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