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Effect of thermo-mechanical cycling on the microstructure and toughness in the weld CGHAZ of a novel high strength low carbon steel
Materials Science & Engineering A 639 (2015) 482–488
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of thermo-mechanical cycling on the microstructure and toughness in the weld CGHAZ of a novel high strength low carbon steel H. Xie a, L.-X. Du a,n, J. Hu a, G.-S. Sun a, H.-Y. Wu a, R.D.K. Misra b,n a
The State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110819, China Laboratory for Excellence in Advanced Steel Research, Center for Structural and Functional Materials Research and Innovation and Department of Metallurgical and Materials Engineering, University of Texas at El Paso, El Paso, TX 79968, USA
b
art ic l e i nf o
a b s t r a c t
Article history: Received 27 April 2015 Received in revised form 9 May 2015 Accepted 11 May 2015 Available online 16 May 2015
We have studied the microstructural evolution in the simulated coarse grain heat affected zone (CGHAZ) of novel low carbon microalloyed steel with yield strength of 1000 MPa using electron microscopy, while the crystallographic characteristics were studied by electron backscatter diffraction (EBSD). The deterioration in low temperature toughness after the simulated welding cycle was attributed to the formation of coarse blocky M–A constituent. However, the lower bainite microstructure of the matrix was beneficial for low temperature impact toughness because of similarity in the crystal structure of variants of lath martensite, which were present in high proportion at the high angle grain boundaries. A high fraction of small M–A constituent also influenced impact toughness. & 2015 Elsevier B.V. All rights reserved.
Keywords: Coarse grain heat affected zone Low carbon steel Ultra-high strength steel Martensite–austenite constituent Toughness
1. Introduction Increasing the strength of a metal requires the use of expensive alloying elements, which in most cases are indispensable. A more direct method is to increase the carbon content, although this promotes the formation of martensite–austenite (M–A) constituent, which can drastically reduce the heat-affected zone (HAZ) toughness [1–3]. Due to this, low carbon microalloyed high strength steels have been widely studied during the past decade. However other weld-related issues must also be taken into consideration because of the possibility of formation of M–A constituent, even at low carbon-content equivalent. In the previous study, we reported that novel Nb–V–Ti and Cr–Mo addition to microalloyed, low carbon content bainitic steel produces an attractive structural material with high strength and excellent low temperature toughness [4]. However, this combination of high strength and high impact toughness may deteriorate after thermal cycling that occurs during welding [5–8], because of austenite grain growth and formation of hard transformation products during cooling. This results in low toughness zones which are susceptible to brittle-fracture within the local brittle zones (LBZs) [9–10]. These LBZs are mainly located within the coarse grained heat affected zone (CGHAZ) and the n
Corresponding authors. E-mail addresses: [email protected] (L.-X. Du), [email protected] (R.D.K. Misra). http://dx.doi.org/10.1016/j.msea.2015.05.033 0921-5093/& 2015 Elsevier B.V. All rights reserved.
intercritically reheated CGHAZ (IC CGHAZ) after welding of the material [11]. Previous studies [2,12] indicated the structure of M–A constituent and its effects on strength and toughness of high strength steel that was welded using various heat inputs. It was postulated that M–A constituent was the main factor that led to deterioration of HAZ toughness. Mohseni et al. [11] studied the X80 pipeline steel and proposed a mechanism by which M–A constituent leads to deterioration of IC CGHAZ toughness. However, some authors have an opposite view point. Li et al. [13] proposed that small grain size and reduced M–A constituent in CGHAZs enhanced toughness. Thus, the effect of M–A constituent on HAZ toughness continues to be unclear. In the present study, the effect of size of M–A constituent in a novel high strength steel is explored by conducting studies on simulated CGHAZ specimens with four different cooling rates. The toughness of the simulated specimens was measured by Charpy impact test, while the influence of welding thermal cycles on microstructural evolution and impact toughness are also studied.
2. Experimental procedure The novel high strength low carbon steel used in this study was a 13-mm thick plate that was thermo-mechanically processed. The chemical composition, mechanical properties, thermo-mechanical processing parameters are summarized in Table 1. The specimens were machined to dimensions of 11 mm 11 mm 55 mm and
H. Xie et al. / Materials Science & Engineering A 639 (2015) 482–488
483
Table 1 Chemical composition and mechanical properties of base metal. Chemical composition (wt%)
Mechanical properties
C
Si
Mn
Nb þV þ Ti
Cr þMo
B
Ceq
Pcm
Rp0.2 (MPa)
Rm (MPa)
A (%)
Toughness (J)
0.08–0.11
0.4
1.0–1.9
0.12
1.08
0.0013
0.72
0.28
1058
1192
10.1
24
Note: Ceq ¼C þ (Mn þSi)/6 þ (Niþ Cu)/15 þ(Cr þ Moþ V)/5; Pcm ¼C þ Si/30þ (Mn þCu þCr)/20 þ Ni/60 þMo/15 þ5B; Impact toughness of 1/4-Charpy size specimen.
1400
1400 10s 30s 60s 150s
1200 1000
10s 30s 60s 150s
1200 1000
800
800
600
600
400
400
200
200 0
0 0
500
1000
1500
2000
2500
3000
3500
0
500
1000
Time(s)
1500
2000
2500
3000
3500
Time(s)
Fig. 1. Welding thermal cycle of the simulated CGHAZ specimens. (a) Actual welding thermal cycle of the simulated specimens and (b) the modeled plots of the simulated specimens.
M
Block
AF
LB
M
Packet
LB
GB
GB LB GB
Fig. 2. Optical micrographs of simulated CGHAZ with different t8/5, 1000 . (a) 10 s, (b) 30 s, (c) 60 s, and (d) 150 s.
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welding thermal simulation was carried out using a thermo-mechanical simulator. The specimens were heated to a peak temperature of 1350 °C at a rate of 100 °C/min and then held there for 1 s, after which they were cooled to 150 °C using different t8/5 time (time to cool from 800 to 500 °C), and then air cooled to room temperature. The selected t8/5 times were 10, 30, 60 and 150 s, to simulate different welding inputs, which were 16.2, 28.0, 39.6 and 62.7 kJ/cm, respectively, that were calculated using the Rykalin mathematical model [14]. The welding simulations were repeated three times for each t8/5 time condition, from which the average impact toughness was obtained. After simulation, the specimens were machined to standard Charpy v-notch specimens and the test was performed at 20 °C using Instron Dynatup 9200 Series drop weight impact tester, consistent with ASTM E23 specification. The fracture surfaces were observed using a FEI Quanta 600 scanning electron microscope (SEM). The simulated specimens were cut near the monitoring thermocouple, metallographically polished, and etched with a 4% nital solution for observing using a Leica DMIRM 2500M optical microscope and Zeiss Ultra 55 SEM. For EBSD studies, the samples were electrolytically polished in a solution of 8% perchloric acid– alcohol solution. EBSD maps were obtained at a step size of 0.1 μm, and analyzed using HKL-Channel 5 software.
3. Results and discussion 3.1. The actual temperature of the simulated CGHAZ specimens Plot of temperature vs. time for the CGHAZ simulation is presented in Fig. 1. The sample temperature was controlled until it dropped to 150 °C at which no phase transformation occurs based on the Rykalin mathematical model. As shown in Fig. 1a and b, the actual temperature agrees with the modeled plots. Variation in cooling times is because of different t8/5 time used to simulate microstructural changes at different cooling rates. 3.2. The microstructure of simulated CGHAZ specimens Fig. 2 shows optical micrographs of simulated CGHAZ specimens at different cooling rates. The figure indicates that microstructural changes dramatically occur with increasing t8/5 cooling time. The microstructure of the specimens subjected to t8/5 of 10 s was characterized by refined martensite (Fig. 2a). The prior austenite grains comprised of large structural units, i.e. packets and blocks [15], having a diameter of 25–40 μm. The average packet width was 10–15 μm, and the block width was 3–5 μm. This is because of growth of austenitic grains at high peak temperature and subsequent fast cooling during simulated welding with a low heat input of 16.2 kJ/cm. The CGHAZ microstructure was observed
LB M M
Block
AF
Packet
3μm
3μm
GB
LB
3μm Fig. 3. SEM micrographs of the simulated CGHAZ with different t8/5, 2000 . (a) 10 s, (b) 30 s, (c) 60 s, and (d) 150 s.
3μm
H. Xie et al. / Materials Science & Engineering A 639 (2015) 482–488
to coarsen on increasing cooling time. At t8/5 time of 30 s (Fig. 2b), the microstructure consisted of acicular ferrite, lath bainite, and martensite. The martensite fraction decreased with increase in heat input, whereas formation of lath bainite and the increase in local C-concentration led to the formation of acicular ferrite within prior austenite. At t8/5 time of 60 s (Fig. 2c), the microstructure was predominantly comprised of lath bainite, with width greater than martensite. Hence, grain size was also increased with increase in heat input [16]. At t8/5 time of 150 s (Fig. 2d), granular bainite and lath bainite microstructure was observed, which was different from other three steels, with proeutectoid ferrite present at the prior austenite grain boundaries, because of low cooling rate used. The M–A constituent is appreciably larger and was 2 μm long and 1 μm wide. As described above, the microstructure was significantly influenced by different heat inputs. Further study of the resulting fine microstructure was performed using SEM and is presented in Fig. 3. At t8/5 of 10 s, it was possible to easily identify a packet in the martensite microstructure containing various blocks (Fig. 3a) and a
block consisting of several laths with straight boundaries. At t8/5 of 30 s, the microstructure comprised of martensite, lath bainite, and acicular ferrite, as shown in Fig. 3b. Because of low cooling rate, the small phase inside the microstructure exhibited slight coarsening and M–A constituent appeared. At t8/5 of 60 s, (Fig. 3c), the microstructure entirely comprised of lath bainite, which was coarser than the martensitic structure, although laths had a crystallographic relationship similar to martensite. At t8/5 of 150 s, the microstructure exhibited necklace-like structure of block M–A constituent present at the grain boundaries. The M–A constituent and typical granular bainite were present inside the prior austenite grain (Fig. 3d). It was also noted that the microstructure was significantly coarse, and the spacing between different phases was appreciably large. 3.3. The crystallographic characteristics of the simulated CGHAZ specimens The crystallographic characteristics of the simulated CGHAZ specimens were analyzed by EBSD and are presented in Fig. 4. The
Selected grain A
Selected grain B
25μm
25μm
25μm
25 μm
{100}
{100}
0
485
0
Fig. 4. Crystallographic characteristics of the simulated CGHAZ specimens analyzed by EBSD: (a) the orientation image map with t8/5 of 30 s, (b) the orientation image map with t8/5 of 60 s, (c) grain boundary misorientation distribution with t8/5 of 30 s, (d) grain boundary misorientation distribution with t8/5 of 60 s, (e) experimental 〈100〉 pole figure containing martensitic variants within the selected austenite grain A, and (f) experimental 〈100〉 pole figure containing bainitic variants within the selected austenite grain B. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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5μm
5μm
5μm Fig. 5. The image quality maps with the austenite distribution of the simulated CGHAZ specimens with different t8/5. (a) 10 s, (b) 60 s, and (c) 150 s.
Total impact energy Crack initiation energy
60
Energy(J)
50
40
30
20
10
0
20
40
60
80
100
120
140
160
t8/5(s) Fig. 6. Total impact energy and crack initiation energy of the simulated CGHAZ specimens.
specimens with t8/5 of 30 s exhibited typical lath martensitic microstructure with various variants, as shown in Fig. 4a. The width of the variants in specimens with t8/5 of 60 s was larger because of lower degree of supercooling. Fig. 4c and d are grain boundary images of the simulated CGHAZ specimens. Large-angle grain boundaries are indicated by black lines, and low-angle grain boundaries are indicated in blue. Because of the high cooling rate, the homogeneity of the large angle grain boundaries in specimens
with t8/5 of 30 s was not as significant as that in specimens with t8/5 of 60 s. Analysis of 〈110〉 pole figures obtained from the selected initial austenite A and B in specimens with t8/5 of 30 and 60 s (Fig. 4e and f), indicated that the orientation relationship of the microstructures was close to Nishiyama–Wassermann than Kurdjumov–Sachs [17,18]. The selected prior austenite A in specimens with t8/5 of 30 s contained nine variants (Fig. 4e), while there are twelve variants in one prior austenite grain B in specimens with t8/5 of 60 s (Fig. 4f). Thus, the crystallographic orientation of specimens with t8/5 of 60 s was similar to specimens with t8/5 of 30 s. Fig. 5 is EBSD image quality maps with distribution of M–A constituent within the simulated CGHAZ specimens. As previously mentioned, the volume and the size of M–A constituents were observed to increase with increasing t8/5 time. However, the austenitic fraction in M–A constituent inside the prior austenite was dominant, whereas the martensitic fraction in M–A constituent dispersed at the rim of prior austenite was considerably large, as shown in Fig. 5b and c. It can be said that the retained austenite relaxes the local stress concentration [19,20]. However, significant fraction of M–A constituent consisted of martensite that was present at the grain boundaries of the prior austenite. The martensite in the M–A constituent contained a high density of dislocations [21], which would severely influence impact toughness. 3.4. The mechanical properties of the simulated CGHAZ specimens Fig. 6 shows the total impact energy and the crack initiation energy of the simulated CGHAZ specimens and indicates that toughness first increases and then decreases with increasing t8/5 time. This is due to the reduction of the brittle martensitic phase at
H. Xie et al. / Materials Science & Engineering A 639 (2015) 482–488
487
Small dimples
Fig. 7. SEM micrographs of fracture surfaces of the simulated CGHAZ specimens. (a) 10 s, (b) 30 s, (c) 60 s, and (d) 150 s.
t8/5 times less than and up to 60 s, whereas the reduced toughness observed on exceeding 60 s is due to the formation of blocky M–A constituents composed of martensitic proportions. It is believed that the appearance of coarse M–A constituents in specimens with cooling time of 150 s was the main reason for the dramatic decreasing in impact toughness of the simulated CGHAZ specimens. The fracture surface of specimens with different t8/5 time is presented in Fig. 7, and the fractographs correspond to the radial regions of specimens. Only highly misoriented boundaries can be transformed to cleavage facet boundaries [3], and it is observed that the cleavage facet size on the specimen fracture surface was relatively large for samples with a cooling time of 10 s. The facet size was also equal to the packet size, as shown in Fig 7a. On increasing t8/5 time, the microstructure of the simulated CGHAZ specimens transformed to lath bainite. Thus, the cleavage facet size was decreased and was separated by boundaries of bainitic lath, as shown in Fig. 7b. It is well known that the formation of fine retained austenitic structure distributed at the grain boundaries of prior austenite grain can effectively release stress concentration [19]. Examination of fractograph of specimen with t8/5 of 60 s indicated that there was some ductile fracture with large number of small dimples (indicated by white arrows in Fig. 7c), which is in good agreement with the study of Lan [3]. Increasing t8/5 time led to the appearance of coarse M–A constituent in the sample, which was detrimental to toughness, as evinced by cleavage planes on the fractograph of the specimen with t8/5 of 150 s (Fig. 7d).
4. Conclusions The change in morphology and mechanical properties of the simulated CGHAZ specimens with different t8/5 time were studied. The major conclusions are summarized as follows: (1) The microstructure of the simulated CGHAZ specimen changed from lath martensite to lath bainite and then to coarse granular bainite with increase in t8/5 time. This led to an increase and then slight decrease in impact energy. (2) The EBSD analysis indicated that lath bainite retains crystallographic orientations similar to lath martensite. The orientations of variants transformed from prior austenite were close to Nishiyama–Wassermann orientation relationship rather than the Kurdjumov–Sachs orientation relationship. Orientation relationship analysis indicated that there was a relationship between the transformed phase and the old phase. The homogeneity of large-angle grain boundaries of lath bainite was better than lath martensite, which benefited the impact toughness of the simulated CGHAZ specimens. (3) The formation of coarse M–A constituent because of low cooling rate is the main reason for the decrease in impact toughness of the simulated CGHAZ specimens. The coarse M–A constituent was nearly fully martensitic, which raises the stress concentration at the rim of the prior austenite and leading to cleavage crack.
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Acknowledgments The authors appreciate the support from the National High Technology Research and Development Program of China (863 Program 2015AA03A501), and the National Natural Science Foundation of China (Grant no. 50527402). RDKM gratefully acknowledges support from the Department of Metallurgical and Materials Engineering, University of Texas at El Paso, USA.
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A. Lambert-Perlade, A.F. Gourgues, A. Pineau, Acta Mater. 52 (2004) 2337–2348. S. Moeinifar, A.H. Kokabi, H.R. Madaah Hosseini, Mater. Des. 32 (2011) 869–876. L.Y. Lan, C.L. Qiu, D.W. Zhao, X.H. Gao, L.X. Du, Mater. Sci. Eng. A 529 (2011) 192–200. H. Xie, L.X. Du, J. Hu, R.D.K. Misra, Mater. Sci. Eng. A 612 (2014) 123–130. C. Thaulow, A.J. Paauw, K. Guttormsen, Int. Inst. Weld. 59 (1987) 266–279. B.C. Kim, S. Lee, N.J. Kim, D.Y. Lee, Metall. Trans. A 22 (1991) 139–149.
[7] A. Lambert-Perlade, A.F. Gourgues, J. Besson, T. Sture, A. Pineau, Metall. Mater. Trans. A 35 (2004) 1039–1053. [8] J. Moon, S. Kim, J. Lee, B. Hwang, C.G. Lee, C. Lee, Mater. Sci. Eng. A 497 (2008) 153–159. [9] C.L. Davis, J.E. King, Metall. Mater. Trans. A 25 (1994) 563–573. [10] A.E. Amer, M.Y. Koo, K.H. Lee, S.H. Kim, S.H. Hong, J. Mater. Sci. 45 (2010) 1248–1254. [11] P. Mohseni, J.K. Solberg, M. Karlsen, O.M. Akselsen, Mater. Sci. Technol. 28 (2012) 1261–1268. [12] S. Moeinifar, A.H. Kokabi, H.R. Madaah Hosseini, Mater. Des. 31 (2010) 2948–2955. [13] C.W. Li, Y. Wang, Y.H. Chen, J. Mater. Sci. 46 (2011) 6424–6431. [14] N.N. Rykalin, Calculation of Heat Processes in Welding, U.S.S.R., Moscow, 1960. [15] S. Morito, X. Huang, T. Furuhara, T. Maki, N. Hansen, Acta Mater. 54 (2006) 5323–5331. [16] S. Kim, Y. Kang, C. Lee, Mater. Sci. Eng. A 559 (2013) 178–186. [17] P.P. Suikkanen, C. Cayron, A.J. DeArdo, L.P. Karjalainen, J. Mater. Sci. Technol. 27 (2011) 920–930. [18] P.P. Suikkanen, C. Cayron, A.J. DeArdo, L.P. Karjalainen, J. Mater. Sci. Technol. 29 (2013) 359–366. [19] C.G. Lee, S.J. Kim, T.H. Lee, S. Lee, Mater. Sci. Eng. A 371 (2004) 16–23. [20] O. Matsumura, Y. Sakuma, Y. Ishii, J. Zhao, ISIJ Int. 32 (1992) 1110–1116. [21] G. Krauss, Mater. Sci. Eng. A 273–275 (1999) 40–57.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of thermo-mechanical cycling on the microstructure and toughness in the weld CGHAZ of a novel high strength low carbon steel H. Xie a, L.-X. Du a,n, J. Hu a, G.-S. Sun a, H.-Y. Wu a, R.D.K. Misra b,n a
The State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110819, China Laboratory for Excellence in Advanced Steel Research, Center for Structural and Functional Materials Research and Innovation and Department of Metallurgical and Materials Engineering, University of Texas at El Paso, El Paso, TX 79968, USA
b
art ic l e i nf o
a b s t r a c t
Article history: Received 27 April 2015 Received in revised form 9 May 2015 Accepted 11 May 2015 Available online 16 May 2015
We have studied the microstructural evolution in the simulated coarse grain heat affected zone (CGHAZ) of novel low carbon microalloyed steel with yield strength of 1000 MPa using electron microscopy, while the crystallographic characteristics were studied by electron backscatter diffraction (EBSD). The deterioration in low temperature toughness after the simulated welding cycle was attributed to the formation of coarse blocky M–A constituent. However, the lower bainite microstructure of the matrix was beneficial for low temperature impact toughness because of similarity in the crystal structure of variants of lath martensite, which were present in high proportion at the high angle grain boundaries. A high fraction of small M–A constituent also influenced impact toughness. & 2015 Elsevier B.V. All rights reserved.
Keywords: Coarse grain heat affected zone Low carbon steel Ultra-high strength steel Martensite–austenite constituent Toughness
1. Introduction Increasing the strength of a metal requires the use of expensive alloying elements, which in most cases are indispensable. A more direct method is to increase the carbon content, although this promotes the formation of martensite–austenite (M–A) constituent, which can drastically reduce the heat-affected zone (HAZ) toughness [1–3]. Due to this, low carbon microalloyed high strength steels have been widely studied during the past decade. However other weld-related issues must also be taken into consideration because of the possibility of formation of M–A constituent, even at low carbon-content equivalent. In the previous study, we reported that novel Nb–V–Ti and Cr–Mo addition to microalloyed, low carbon content bainitic steel produces an attractive structural material with high strength and excellent low temperature toughness [4]. However, this combination of high strength and high impact toughness may deteriorate after thermal cycling that occurs during welding [5–8], because of austenite grain growth and formation of hard transformation products during cooling. This results in low toughness zones which are susceptible to brittle-fracture within the local brittle zones (LBZs) [9–10]. These LBZs are mainly located within the coarse grained heat affected zone (CGHAZ) and the n
Corresponding authors. E-mail addresses: [email protected] (L.-X. Du), [email protected] (R.D.K. Misra). http://dx.doi.org/10.1016/j.msea.2015.05.033 0921-5093/& 2015 Elsevier B.V. All rights reserved.
intercritically reheated CGHAZ (IC CGHAZ) after welding of the material [11]. Previous studies [2,12] indicated the structure of M–A constituent and its effects on strength and toughness of high strength steel that was welded using various heat inputs. It was postulated that M–A constituent was the main factor that led to deterioration of HAZ toughness. Mohseni et al. [11] studied the X80 pipeline steel and proposed a mechanism by which M–A constituent leads to deterioration of IC CGHAZ toughness. However, some authors have an opposite view point. Li et al. [13] proposed that small grain size and reduced M–A constituent in CGHAZs enhanced toughness. Thus, the effect of M–A constituent on HAZ toughness continues to be unclear. In the present study, the effect of size of M–A constituent in a novel high strength steel is explored by conducting studies on simulated CGHAZ specimens with four different cooling rates. The toughness of the simulated specimens was measured by Charpy impact test, while the influence of welding thermal cycles on microstructural evolution and impact toughness are also studied.
2. Experimental procedure The novel high strength low carbon steel used in this study was a 13-mm thick plate that was thermo-mechanically processed. The chemical composition, mechanical properties, thermo-mechanical processing parameters are summarized in Table 1. The specimens were machined to dimensions of 11 mm 11 mm 55 mm and
H. Xie et al. / Materials Science & Engineering A 639 (2015) 482–488
483
Table 1 Chemical composition and mechanical properties of base metal. Chemical composition (wt%)
Mechanical properties
C
Si
Mn
Nb þV þ Ti
Cr þMo
B
Ceq
Pcm
Rp0.2 (MPa)
Rm (MPa)
A (%)
Toughness (J)
0.08–0.11
0.4
1.0–1.9
0.12
1.08
0.0013
0.72
0.28
1058
1192
10.1
24
Note: Ceq ¼C þ (Mn þSi)/6 þ (Niþ Cu)/15 þ(Cr þ Moþ V)/5; Pcm ¼C þ Si/30þ (Mn þCu þCr)/20 þ Ni/60 þMo/15 þ5B; Impact toughness of 1/4-Charpy size specimen.
1400
1400 10s 30s 60s 150s
1200 1000
10s 30s 60s 150s
1200 1000
800
800
600
600
400
400
200
200 0
0 0
500
1000
1500
2000
2500
3000
3500
0
500
1000
Time(s)
1500
2000
2500
3000
3500
Time(s)
Fig. 1. Welding thermal cycle of the simulated CGHAZ specimens. (a) Actual welding thermal cycle of the simulated specimens and (b) the modeled plots of the simulated specimens.
M
Block
AF
LB
M
Packet
LB
GB
GB LB GB
Fig. 2. Optical micrographs of simulated CGHAZ with different t8/5, 1000 . (a) 10 s, (b) 30 s, (c) 60 s, and (d) 150 s.
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welding thermal simulation was carried out using a thermo-mechanical simulator. The specimens were heated to a peak temperature of 1350 °C at a rate of 100 °C/min and then held there for 1 s, after which they were cooled to 150 °C using different t8/5 time (time to cool from 800 to 500 °C), and then air cooled to room temperature. The selected t8/5 times were 10, 30, 60 and 150 s, to simulate different welding inputs, which were 16.2, 28.0, 39.6 and 62.7 kJ/cm, respectively, that were calculated using the Rykalin mathematical model [14]. The welding simulations were repeated three times for each t8/5 time condition, from which the average impact toughness was obtained. After simulation, the specimens were machined to standard Charpy v-notch specimens and the test was performed at 20 °C using Instron Dynatup 9200 Series drop weight impact tester, consistent with ASTM E23 specification. The fracture surfaces were observed using a FEI Quanta 600 scanning electron microscope (SEM). The simulated specimens were cut near the monitoring thermocouple, metallographically polished, and etched with a 4% nital solution for observing using a Leica DMIRM 2500M optical microscope and Zeiss Ultra 55 SEM. For EBSD studies, the samples were electrolytically polished in a solution of 8% perchloric acid– alcohol solution. EBSD maps were obtained at a step size of 0.1 μm, and analyzed using HKL-Channel 5 software.
3. Results and discussion 3.1. The actual temperature of the simulated CGHAZ specimens Plot of temperature vs. time for the CGHAZ simulation is presented in Fig. 1. The sample temperature was controlled until it dropped to 150 °C at which no phase transformation occurs based on the Rykalin mathematical model. As shown in Fig. 1a and b, the actual temperature agrees with the modeled plots. Variation in cooling times is because of different t8/5 time used to simulate microstructural changes at different cooling rates. 3.2. The microstructure of simulated CGHAZ specimens Fig. 2 shows optical micrographs of simulated CGHAZ specimens at different cooling rates. The figure indicates that microstructural changes dramatically occur with increasing t8/5 cooling time. The microstructure of the specimens subjected to t8/5 of 10 s was characterized by refined martensite (Fig. 2a). The prior austenite grains comprised of large structural units, i.e. packets and blocks [15], having a diameter of 25–40 μm. The average packet width was 10–15 μm, and the block width was 3–5 μm. This is because of growth of austenitic grains at high peak temperature and subsequent fast cooling during simulated welding with a low heat input of 16.2 kJ/cm. The CGHAZ microstructure was observed
LB M M
Block
AF
Packet
3μm
3μm
GB
LB
3μm Fig. 3. SEM micrographs of the simulated CGHAZ with different t8/5, 2000 . (a) 10 s, (b) 30 s, (c) 60 s, and (d) 150 s.
3μm
H. Xie et al. / Materials Science & Engineering A 639 (2015) 482–488
to coarsen on increasing cooling time. At t8/5 time of 30 s (Fig. 2b), the microstructure consisted of acicular ferrite, lath bainite, and martensite. The martensite fraction decreased with increase in heat input, whereas formation of lath bainite and the increase in local C-concentration led to the formation of acicular ferrite within prior austenite. At t8/5 time of 60 s (Fig. 2c), the microstructure was predominantly comprised of lath bainite, with width greater than martensite. Hence, grain size was also increased with increase in heat input [16]. At t8/5 time of 150 s (Fig. 2d), granular bainite and lath bainite microstructure was observed, which was different from other three steels, with proeutectoid ferrite present at the prior austenite grain boundaries, because of low cooling rate used. The M–A constituent is appreciably larger and was 2 μm long and 1 μm wide. As described above, the microstructure was significantly influenced by different heat inputs. Further study of the resulting fine microstructure was performed using SEM and is presented in Fig. 3. At t8/5 of 10 s, it was possible to easily identify a packet in the martensite microstructure containing various blocks (Fig. 3a) and a
block consisting of several laths with straight boundaries. At t8/5 of 30 s, the microstructure comprised of martensite, lath bainite, and acicular ferrite, as shown in Fig. 3b. Because of low cooling rate, the small phase inside the microstructure exhibited slight coarsening and M–A constituent appeared. At t8/5 of 60 s, (Fig. 3c), the microstructure entirely comprised of lath bainite, which was coarser than the martensitic structure, although laths had a crystallographic relationship similar to martensite. At t8/5 of 150 s, the microstructure exhibited necklace-like structure of block M–A constituent present at the grain boundaries. The M–A constituent and typical granular bainite were present inside the prior austenite grain (Fig. 3d). It was also noted that the microstructure was significantly coarse, and the spacing between different phases was appreciably large. 3.3. The crystallographic characteristics of the simulated CGHAZ specimens The crystallographic characteristics of the simulated CGHAZ specimens were analyzed by EBSD and are presented in Fig. 4. The
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Fig. 4. Crystallographic characteristics of the simulated CGHAZ specimens analyzed by EBSD: (a) the orientation image map with t8/5 of 30 s, (b) the orientation image map with t8/5 of 60 s, (c) grain boundary misorientation distribution with t8/5 of 30 s, (d) grain boundary misorientation distribution with t8/5 of 60 s, (e) experimental 〈100〉 pole figure containing martensitic variants within the selected austenite grain A, and (f) experimental 〈100〉 pole figure containing bainitic variants within the selected austenite grain B. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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5μm
5μm
5μm Fig. 5. The image quality maps with the austenite distribution of the simulated CGHAZ specimens with different t8/5. (a) 10 s, (b) 60 s, and (c) 150 s.
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specimens with t8/5 of 30 s exhibited typical lath martensitic microstructure with various variants, as shown in Fig. 4a. The width of the variants in specimens with t8/5 of 60 s was larger because of lower degree of supercooling. Fig. 4c and d are grain boundary images of the simulated CGHAZ specimens. Large-angle grain boundaries are indicated by black lines, and low-angle grain boundaries are indicated in blue. Because of the high cooling rate, the homogeneity of the large angle grain boundaries in specimens
with t8/5 of 30 s was not as significant as that in specimens with t8/5 of 60 s. Analysis of 〈110〉 pole figures obtained from the selected initial austenite A and B in specimens with t8/5 of 30 and 60 s (Fig. 4e and f), indicated that the orientation relationship of the microstructures was close to Nishiyama–Wassermann than Kurdjumov–Sachs [17,18]. The selected prior austenite A in specimens with t8/5 of 30 s contained nine variants (Fig. 4e), while there are twelve variants in one prior austenite grain B in specimens with t8/5 of 60 s (Fig. 4f). Thus, the crystallographic orientation of specimens with t8/5 of 60 s was similar to specimens with t8/5 of 30 s. Fig. 5 is EBSD image quality maps with distribution of M–A constituent within the simulated CGHAZ specimens. As previously mentioned, the volume and the size of M–A constituents were observed to increase with increasing t8/5 time. However, the austenitic fraction in M–A constituent inside the prior austenite was dominant, whereas the martensitic fraction in M–A constituent dispersed at the rim of prior austenite was considerably large, as shown in Fig. 5b and c. It can be said that the retained austenite relaxes the local stress concentration [19,20]. However, significant fraction of M–A constituent consisted of martensite that was present at the grain boundaries of the prior austenite. The martensite in the M–A constituent contained a high density of dislocations [21], which would severely influence impact toughness. 3.4. The mechanical properties of the simulated CGHAZ specimens Fig. 6 shows the total impact energy and the crack initiation energy of the simulated CGHAZ specimens and indicates that toughness first increases and then decreases with increasing t8/5 time. This is due to the reduction of the brittle martensitic phase at
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Small dimples
Fig. 7. SEM micrographs of fracture surfaces of the simulated CGHAZ specimens. (a) 10 s, (b) 30 s, (c) 60 s, and (d) 150 s.
t8/5 times less than and up to 60 s, whereas the reduced toughness observed on exceeding 60 s is due to the formation of blocky M–A constituents composed of martensitic proportions. It is believed that the appearance of coarse M–A constituents in specimens with cooling time of 150 s was the main reason for the dramatic decreasing in impact toughness of the simulated CGHAZ specimens. The fracture surface of specimens with different t8/5 time is presented in Fig. 7, and the fractographs correspond to the radial regions of specimens. Only highly misoriented boundaries can be transformed to cleavage facet boundaries [3], and it is observed that the cleavage facet size on the specimen fracture surface was relatively large for samples with a cooling time of 10 s. The facet size was also equal to the packet size, as shown in Fig 7a. On increasing t8/5 time, the microstructure of the simulated CGHAZ specimens transformed to lath bainite. Thus, the cleavage facet size was decreased and was separated by boundaries of bainitic lath, as shown in Fig. 7b. It is well known that the formation of fine retained austenitic structure distributed at the grain boundaries of prior austenite grain can effectively release stress concentration [19]. Examination of fractograph of specimen with t8/5 of 60 s indicated that there was some ductile fracture with large number of small dimples (indicated by white arrows in Fig. 7c), which is in good agreement with the study of Lan [3]. Increasing t8/5 time led to the appearance of coarse M–A constituent in the sample, which was detrimental to toughness, as evinced by cleavage planes on the fractograph of the specimen with t8/5 of 150 s (Fig. 7d).
4. Conclusions The change in morphology and mechanical properties of the simulated CGHAZ specimens with different t8/5 time were studied. The major conclusions are summarized as follows: (1) The microstructure of the simulated CGHAZ specimen changed from lath martensite to lath bainite and then to coarse granular bainite with increase in t8/5 time. This led to an increase and then slight decrease in impact energy. (2) The EBSD analysis indicated that lath bainite retains crystallographic orientations similar to lath martensite. The orientations of variants transformed from prior austenite were close to Nishiyama–Wassermann orientation relationship rather than the Kurdjumov–Sachs orientation relationship. Orientation relationship analysis indicated that there was a relationship between the transformed phase and the old phase. The homogeneity of large-angle grain boundaries of lath bainite was better than lath martensite, which benefited the impact toughness of the simulated CGHAZ specimens. (3) The formation of coarse M–A constituent because of low cooling rate is the main reason for the decrease in impact toughness of the simulated CGHAZ specimens. The coarse M–A constituent was nearly fully martensitic, which raises the stress concentration at the rim of the prior austenite and leading to cleavage crack.
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Acknowledgments The authors appreciate the support from the National High Technology Research and Development Program of China (863 Program 2015AA03A501), and the National Natural Science Foundation of China (Grant no. 50527402). RDKM gratefully acknowledges support from the Department of Metallurgical and Materials Engineering, University of Texas at El Paso, USA.
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