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Dynamic fracture behavior of low carbon bainitic steel after different welding thermal cycles
Engineering Fracture Mechanics 220 (2019) 106653
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Dynamic fracture behavior of low carbon bainitic steel after different welding thermal cycles
T
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Liangyun Lana,b, , Yiting Zhanga, Xiangwei Konga,b a
School of Mechanical Engineering and Automation, Northeastern University, Shenyang 110819, China Key Laboratory of Vibration and Control of Aero-Propulsion Systems, Ministry of Education of China, Northeastern University, Shenyang 110819, China
b
A R T IC LE I N F O
ABS TRA CT
Keywords: Dynamic fracture toughness Low carbon steel Stretch zone width Crack initiation point
Dynamic fracture toughness is an important parameter to assess the fracture resistance at crack initiation although the actual crack initiation point is hard to be precisely detected. In this work, three different heat treatment microstructures (i.e., hot-rolled, welding thermal simulation coarse grained HAZ (CGHAZ) and intercritically reheated CGHAZ microstructures) were studied by instrumented Charpy impact tester. The dynamic fracture toughness was obtained based on impact load-deflection curves using various estimation methods. The dynamic fracture toughness decreased with the applied welding thermal cycles irrespective of the estimation methods. The stretch zone width proved that the fracture toughness is most acceptable when the crack initiation point is considered as the middle point between the general yield load and maximum load in the load-deflection curves.
1. Introduction Welding is one of the most common methods in industrial practice for joining components because of low costs and high efficiency. However, it also yields many defects, e.g., heterogeneous microstructure, local brittle zone, cold cracks and residual stress. These defects threat the integrity and reliability of weldment structure [1]. Thus, the attention from metallurgists has been always paid to the reasons that cause the deterioration in mechanical properties of welded joints. Neves and Loureiro [2] showed that a decrease in toughness was observed in the weld metal and heat-affected zone (HAZ) with increasing heat input as a result of the increasing size of the local brittle zone sampled by fatigue crack front. Yang et al. [3] found that the fracture toughness of fusion zone is poorest in the X80 pipeline weldment and the hard and brittle martensite-austenite (MA) constituents may play roles of the fracture origins. Linear elastic fracture mechanics based plane strain fracture toughness (KIC) approach or Elastic-plastic fracture mechanics based J-integral approach is recognized as an applicable method to carry out structural integrity assessment [1–3]. Determinations of these static fracture toughness parameters, however, require very strict and somewhat complex procedures [4] such as large and thick samples, sharp fatigue pre-crack with reasonable initial length. Particularly for the HAZ with size limits and irregular shape, the sample preparation is hard to meet these requirements. Charpy impact test is also widely applied to reveal the impact fracture behaviors of materials in industry due to the simple, rapid and inexpensive testing procedure. However, from the fracture mechanics point of view, the impact absorbed energy itself cannot be directly used for quantitative safety assessment [5]. Because impact specimen only has a Charpy V or U notch, the total impact energy should be composed of the crack initiation and
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Corresponding author at: School of Mechanical Engineering and Automation, Northeastern University, Shenyang 110819, China. E-mail addresses: [email protected] (L. Lan), [email protected] (X. Kong).
https://doi.org/10.1016/j.engfracmech.2019.106653 Received 7 April 2019; Received in revised form 29 August 2019; Accepted 30 August 2019 Available online 30 August 2019 0013-7944/ © 2019 Elsevier Ltd. All rights reserved.
Engineering Fracture Mechanics 220 (2019) 106653
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propagation energies and the ratio of them mainly depends on the stress-strain state that the samples undergo. Besides the mathematical relationships established between the fracture toughness and Charpy impact energy [6,7], dynamic fracture toughness that is derived from the impact load-deflection curve is another important parameter to assess the fracture resistance at crack initiation [8,9]. However, the crack initiation point is hard to be detected exactly even using instrumented impact tester. In general, the peak load that occurs in the load-deflection curve is approximately regarded as the onset point of crack propagation [10]. It is noted that this method intentionally improves the crack initiation energy, especially under the condition of ductile fracture. Several other methods were proposed in literature to ascertain the crack initiation point, such as compliance changing rate method [8] and averaging method of yield and peak loads [9]. As a result, the dynamic fracture toughness (JId) can be estimated with these methods. Zhu et al. [11] considered that the JId is acceptable when the crack initiation point is considered to occur at a load of (Pm + Py)/2, where Pm and Py is maximum and general yield load, respectively. Moitra et al. [12] showed that the JId follows a power law relationship with the inter-particle spacing for different regions in the HAZ of a nuclear grade steel. Although the dynamic fracture toughness reflects the resistance of materials to macro-crack initiation at the impact load condition, majority of studies only focused on the effect of microstructure on Charpy impact energy rather than dynamic fracture toughness [13–16]. As we known, the Charpy impact energy cannot be directly used for the structural integrity assessment. Meanwhile, the total absorbed energy does not characterize the resistance of materials to crack initiation or propagation at a certain stage because different stages of fracture have some different physical phenomena [17]. In this work, three different heat-treated microstructures, i.e., as-rolled state, welding thermal simulated-CGHAZ (coarse grained HAZ) and ICCGHAZ (intercritically reheated CGHAZ), were employed to study the relationship between microstructure and dynamic fracture toughness. In combination with the morphology of fracture surface, the underlying mechanism that deteriorates the dynamic fracture toughness of the ICCGHAZ was explored.
2. Experimental methods The chemical composition of studied low carbon steel was 0.053 C, 0.25 Si, 1.68 Mn, 0.45 (Cr + Mo), 0.086 (Nb + V + Ti), 0.53 (Cu + Ni), 0.0012 B (Wt.%) balanced by Fe, similar to the composite of X80 pipeline steel [3]. Hot rolled steel plates with a thickness of 13 mm were produced using thermo-mechanically controlled processing (TMCP) in our laboratory and the detailed TMCP parameters as well as as-rolled mechanical properties were reported in Ref. [18]. Here, some important curves were reused to further confirm dynamic fracture parameters and these values were compared with as-welded dynamic fracture parameters. Specimens with a dimension of 11 × 11 × 55 mm were cut from the hot rolled steel plates with their longest dimension perpendicular to rolling direction. Single and double-pass welding thermal cycles were simulated to obtain CGHAZ and ICCGHAZ microstructures, respectively, employing a resistance heating thermo-simulator. For the first welding thermal cycle, the peak temperature was set at 1350 °C to make the austenite grain coarsening; while for the second cycle, the peak temperature was set at 740 °C that is just higher than the α to γ phase transformation temperature (i.e., Ac1 ~ 710 °C). The cooling rate for both cycles was adopted as a cooling time (t8/5 is the elapsed time when the temperature dropped from 800 °C to 500 °C) of 30 s, which is supposed to simulate submerged arc welding with the heat input of about 38 kJ/mm. Fig. 1 shows the measured temperature curve for double-pass welding thermal cycle. Each of simulation processes was conducted at least three times using the fresh specimens. After the welding simulation, all specimens were re-machined into the standard size (10 × 10 × 55 mm) of Charpy impact samples and a 2-mm deep V-notch was machined at the region of interest. The impact testing was performed at −20 °C in an instrumented drop weight impact tester equipped with an oscilloscope. The actual load–deflection curves were recorded with the oscilloscope. The microstructures as well as the morphology of second cracks underneath the fracture surface were examined by a Leica DMIRM optical microscopy after the metallographic specimens were prepared by conventional methods, i.e. manual and mechanical polishing, 3% nital etching and hot air drying. Microhardness measurement was conducted using a FM700 hardnesstesting machine, employed a 0.49 N load. At least 7 indenters were detected for each kind of heat treated microstructures and average
Fig. 1. Measured temperature profile for simulating double-pass welding thermal cycles (P1 and P2: two peak temperatures). 2
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Fig. 2. Raw data impact load- deflection curves for the specimens with different heat treatment states (a) as-rolled state, (b) CGHAZ, (c) ICCGHAZ (three curves given in each condition represent three independent results).
values were reported below. A Zeiss Ultra 55 field emission scanning electron microscope (SEM) was used to observe the morphology of fracture surfaces of broken Charpy specimens.
3. Results and discussion 3.1. Impact load-deflection curves Fig. 2 represents raw impact load-deflection curves of the specimens after subjected to different heat treatments. These curves can be sub-divided into five different stages, i.e., elastic and plastic deformation, ductile propagation stage, brittle (unstable) fracture stage, and ductile fracture stage [10,19], although not all of these stages exist on each curve. For the hot-rolled specimens, their curves always exhibited a relatively stable fracture stage, implying that the ductile crack propagation occurs [20]. Occasionally, some of them also have a very short stage of unstable crack extension (as arrowed in Fig. 2a). In contrast, all of the welding thermally treated specimens exhibited predominantly brittle unstable fracture behaviors since the loads abruptly dropped to zero immediately after the peak load (Fig. 2b and c). This behavior not only leads to a huge reduction in sample’s deflection but also makes the phenomenon of load fluctuations more obvious. Furthermore, the double-pass welding thermal cycle (Fig. 2c) seems to decrease the peak loads as compared to Fig. 2b. To determine the dynamic fracture parameters accurately, these oscillations should be eliminated through data smoothing.
3.2. Dynamic fracture toughness Fig. 3a shows an example of the raw data curve and corresponding smoothened curve. Some dynamic fracture parameters such as Py, Pm and corresponding absorbed energy can be directly obtained from the smoothing curve, as marked with symbols. Fig. 3b represents one group of typical smooth curves for different heat treated specimens. The general yield stress (σy) for full size CVN specimens can be estimated from the Py using the following equation [8]: 3
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Fig. 3. (a) Comparison of raw curve and smoothing curve and the determination of dynamic fracture parameter, (b) one typical group of smooth curves for different heat treated specimens.
σy =
2.99·Py·W B·(W − a0)2
(1)
where B is the sample thickness, W is the sample width and a0 is the initial crack length. For standard Charpy V notch specimens, the B = W = 10 mm, a0/W = 0.2. Based on elastic-plastic fracture mechanics, the JId can be evaluated from the energy absorbed up to the crack initiation point as [8,12],
JId =
ηEa B (W − a0)
(2)
Here, Ea is the absorbed energy up to the crack initiation point and η = 1.45 for standard Charpy V notch specimens. Therefore, the most important procedure to obtain the JId is how to determine the crack initiation point. For simplicity, the crack initiation point is assumed to be at the maximum load point even with an obvious plastic deformation (i.e., Ea = Ei) [21]. Kobayashi [8] developed the compliance changing rate method to determine the crack initiation point and found that the ratio of initiation to maximum load energies is about 0.8 (i.e., Ea/Ei = 0.8). This result is verified by Tronskar et al. [22] using a laser interferometer to measure displacement. Ghoneim and Hammad [9] proposed that the crack initiation is more reasonably taken as at the load point of (Py + Pm)/2 because the crack initiation often occurs above the yield load but prior to maximum load. Zhu et al. [11] and Puppala et al. [23] proved that the JId integral estimated by this method largely agrees with the experimental results. Norris [24] reported the crack initiates at the time of ti/tmax = 0.4, where ti is the time to crack initiation; tmax is the time to reach the maximum load, based on a comparison of finite element results with experimental data. Moitra et al. [12] confirmed that this method gives more conservative estimation of JId than the averaging method of yield and maximum loads. When the Py is almost equal to Pm (e.g. the ICCGHAZ curve in Fig. 3b), this means that the testing temperature approaches to the nil ductility temperature of the specimen. Thus, the fracture stress (σf) can be approximately evaluated from the σy using the following expression:
σf = Cf σy
(3)
where the stress intensification factor Cf is taken as 2.57 for the full size CVN specimens [21]. The details of dynamic fracture parameters determined from curves are shown in Table 1 and the JId values calculated with various methods are presented in Fig. 4, where JId1 represents that the crack initiation occurs at maximum load point, JId2 crack initiation at the potential energy 0.8 times Ei, JId3 crack initiation at the load point of (Py + Pm)/2, and JId4 crack initiation at the time of ti/tmax = 0.4. It can be seen that for these different methods the crack initiation points determined are not the same, which leads to the fact that the dynamic fracture toughness is quite different, especially for the hot rolled specimen. Obviously, the maximum load point represents an over-estimated JId value because it may contain some stable ductile crack propagations after plastic deformation, which Table 1 Average values of dynamic fracture parameters for all impact specimens. Specimen
Py (kN)
Pm (kN)
Ei (J)
Ep (J)
σy (MPa)
σf (MPa)
Hot rolled CGHAZ ICCGHAZ
17.8 17.6 15.6
22.8 22.3 16.8
65 39 8
85 11 2.9
831 822 696
– 2112 1788
4
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Fig. 4. The calculated dynamic fracture toughness with various methods.
will be verified with SEM observation later. The crack initiation point at the time of ti/tmax = 0.4 gives the most conservative JId values as compared to other methods. However, for the ICCGHAZ, the crack initiation point determined with this method is even much lower than the general yield point. It seems to be unreasonable since the crack initiation is hard to occur at so lower stress condition. Compared with the fracture toughness (256–438 kJ/m2 at different regions for X80 pipeline welded joint) obtained by Yang et al.[3], the JId3 value is most acceptable to characterize the resistance of different heat-treated specimens to cracking. In spite of the estimation methods used, the JId always decreases in order of Hot-rolled, CGHAZ, and then ICCGHAZ. This means that the weld thermal cycle deteriorates the fracture toughness. Meanwhile, it can be found that the ratio of Charpy impact energy to JId gradually decreases due to the effect of weld thermal cycles. That is, using the JId3 values the ratio is 0.35 for hot-rolled specimens, 0.18 for CGHAZ, and 0.09 for ICCGHAZ. These phenomena are in good agreement with the results by Lambert-Perlade et al. [24], who considered that the microstructural evolution under the welding thermal cycles is responsible for the deterioration of fracture toughness. 3.3. Microstructures The microstructures with different heat treatments are shown in Fig. 5. The hot rolled specimens have a mixed degenerate upper bainite and lath bainite. The pancake-shape prior austenite grain boundaries can be observed clearly (arrowed in Fig. 5a) and their thickness measured by line intercept method is ~10.4 um. After the specimens were subject to the welding thermal cycles, these pancaked prior austenite grain boundaries disappear and they become near polygonal-shape (arrowed in Fig. 5b) with a coarse average grain size of ~86 um, although the main microstructure is still transformed into upper bainite (or bainitic ferrite). Each coarse prior austenite grain seems to be partitioned into several bainitic packets because the packet morphology is described by the paralleling distribution of carbides and stringer-like MA constituents inside each a bainitic packet [25]. When the specimens were further subjected to the intercritically reheated thermal cycles, the matrix microstructure does not show a notable change except for slightly tempering effect (Fig. 5c) as compared to Fig. 5b. However, it is evident that the prior austenite grain boundaries are decorated by numerous black particles that are considered to be massive MA constituents (as arrowed) and their maximum width ranges from ~3 to 5.6 μm measured at different views of field. These islands are directly derived from the newly formed austenite during the second welding thermal cycles because the peak temperature is just slightly higher than Ac1. And they inherently possess much higher hardness and brittleness than the ferritic matrix due to higher amount of carbon atoms enriched on them [24–26]. Although some difference exists in microstructural morphology for these three different heat treated specimens, the microhardness does not exhibit an obvious change. The average values in HV are 263 ± 12, 270 ± 10, and 268 ± 18 for the hot rolled, CGHAZ and ICCGHAZ specimens, respectively. 3.4. SEM fractography and its relation to fracture toughness The morphology of fracture surface not only reveals the micro-mechanism of fracture but also verifies the mechanical properties by measuring some characteristic parameters, e.g., stretch zone width and unit cleavage facet size [11,24,27–29]. Fig. 6a gives a typical example of macrograph for the fracture surface of as-rolled specimen. The lateral expansion event can be found and the fracture surface is macroscopically rough with the presence of unstable fracture region in the center of the surface (as cycled region). The fibrous stretch zone with a width of ~2 mm appears between the V-notch and unstable fracture region, as marked with double arrows. These two regions correspond to two different fracture mechanisms, i.e., ductile and cleavage fracture shown in Fig. 6c and b, respectively. In combination with Fig. 2a, the change of fracture mechanism from ductile fracture to cleavage fracture is mainly responsible for a short stage of fast decreasing load (arrowed) because the cleavage fracture belongs to a typical low stress brittle fracture [20]. Meanwhile, it also implies that the ductile propagation stage corresponds to the fibrous stretch zone prior to unstable 5
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Fig. 5. Microstructure at different heat treatment states (a) hot-rolled state (RD: rolling direction), (b) CGHAZ, (c) ICCGHAZ.
fracture region. However, the initiation point of ductile macrocrack still cannot be fixed before or after the peak load. In addition, these fine cleavage facets with elongated shape along the rolling direction (Fig. 6b) reflect the crystallographic orientation distribution of microstructure because only high-angle boundaries can effectively arrest the propagation of cleavage cracks [15,24]. Fig. 7 shows the fracture surface of the CGHAZ specimen. From the macrograph (Fig. 7a), the fracture surface is rather flat compared to Fig. 6a and the phenomenon of lateral expansion does not exist. Based on macro-fracture patterns the entire flat fracture seems to initiate at the location near to the notch, as marked with a black grid. The magnified morphology of this square region is shown in Fig. 7b. It is interesting to find that the real initiation site is derived from a coarse TiN precipitate with the size of about 2.7 um (shown in the inset) that is located inside a prior austenite grain. The dotted lines and arrows indicate the propagation paths and directions of cleavage fracture. This precipitate is likely to be cracked by itself to stimulate the initiation of microcrack as the tessellated residual stresses between the TiN and the matrix increase the cleavage-initiating potency of TiN inclusions [30,31]. More interestingly, a very narrow ligament (the width ~0.182 mm) can be found at the bottom of Fig. 7a and it is separated from the cleavage fracture region by a straight line (arrowed in Fig. 7b). This means that the CGHAZ specimen first fractured with ductile tearing for a very short while, and then the cleavage initiation was triggered because the stress concentration region enlarges with the growth of ductile crack [5]. Here, the cleavage initiation site is ~0.8 mm away from the blunt notch. As the cleavage cracks propagate very fast, the front of ductile cracks encounters the cleavage cracks at the same time, which forms an obvious boundary between them. In combination with Fig. 2b, it can be inferred that the ductile tearing must occur to before the peak load because no ductile propagation stage appears after the peak load. That is to say, the increasing load suddenly dropped to zero immediately after the occurrence of cleavage fracture. Under a higher magnification, the unit cleavage facets can be largely identified according to the traces of ridges as delineated in Fig. 7c because these ridges formed are to compensate the local misorientation between two neighboring cleavage facets [17]. The reinitiation site for each facet is likely located at the interior of the grains based on the river-pattern, and the size of unit cleavage facets seems to be comparable to the prior austenite grain size of the CGHAZ (Fig. 5b). For the ICCGHAZ specimen, the morphology of fracture surface is characterized by coarse cleavage facets without obvious ductile tearing region (Fig. 8a). Meanwhile, it can be seen that most of coarse cleavage facets are re-initiated attached to prior austenite grain boundaries as the ridges of cleavage facet are converged at the grain boundary, as arrowed in Fig. 8b. This implies that the microstructural feature at the prior austenite grain boundaries for the ICCGHAZ is in favor of assisting the initiation of microcracks. 6
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Fig. 6. Fracture surface of as-rolled specimens: (a) macrograph, (b) morphology of cleavage fracture at the center of macro- fracture surface and RD indicates rolling direction, (c) morphology of ductile fracture near to the V-notch position.
Davis and King [32] showed that the brittle fracture initiation at prior austenite grain boundaries is attributed to the fact that debonding of these necklace-type MA constituents with matrix microstructure is aggravated under the overlap of residual phase transformation stresses and the stress concentrations. Mohseni et al. [26] confirmed that the initiation sites of cleavage facets are characterized by twinned martensite and argued that the initiation site is hard to be identified due to the divergence of river pattern in the lath-like microstructure. Here, it is acceptable that some particles located at the initiation site of cleavage facets (arrowed in Fig. 8a) are considered to be the MA constituents since they have a similar morphology, size and location to the MA constituents appeared in Fig. 5c. The morphology of secondary cleavage cracks underneath the fracture surface can further illustrate the initiation sites of cleavage crack because the stress state that they undergo may be similar to the main fracture surface. Fig. 8c shows a portion of cross-sectional microstructure in the vicinity of the fracture surface. The arrows in Fig. 8c represent the unit crack paths during main crack propagation. The unit path is defined as the region in which the crack propagates almost in a straight line [29]. It is closely related to the size of cleavage facets. Evidently, most prior austenite grain boundaries inhibit the linear propagation of main cracks effectively although the cracks, sometimes, exhibit some deviations inside the prior austenite grains. Many secondary cracks can be found in the same or nearby prior austenite grains of the main cracks, as marked with a square grid in Fig. 8c. This squared region is magnified in Fig. 8d to show the detailed morphology of secondary cracks. They are approximately parallel to the main fracture surface and most of them seem to initiate individually at massive MA constituents along the prior austenite grain boundaries, as marked with arrows. Especially for these small microcracks, they seem just nucleate on the MA constituents and do not further propagate probably due to fast decreasing stress. Thus, these blocky MA constituents along the prior austenite grain boundaries indeed act as potential crack initiation sites in the ICCGHAZ. Because the ICCGHAZ sample almost exhibits complete brittle fracture, the fracture stress calculated from the general yield load 7
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Fig. 7. Fracture surface of the CGHAZ specimens: (a) macrograph, (b) the initiation of cleavage fracture at the TiN precipitate near to the V notch position, (c) coarse cleavage facets.
(Table 1) should be satisfied with the formation of critical Griffith microcrack size according to the classic Griffith theory: 1/2 π Eγp ⎞ σc = ⎜⎛ ⎟ 2 ⎝ (1 − ν ) d ⎠
(4)
where E, ν, and γp are the material’s parameters, i.e., Young’s modulus, Poisson’s ratio and the effective surface energy of microcrack, respectively. Using the same values given in Ref. [31], the critical Griffith microcrack size is ~ 3.2 μm at the maximum fracture stress. This critical size agrees well with the width size of MA constiutents formed in the ICCGHAZ mentioned above. In contrast, coarse precipitates such as TiN in the CGHAZ, no doubt, acts as a trigger for cleavage fracture (Fig. 7b). It can be calculated that the critical fracture stress needed is ~1938 MPa as the TiN size measured on the SEM image is ~2.7 μm. This required critical stress is lower than the calculated fracture stress (Table 1) in the front of ductile cracks. It is acceptable because the CGHAZ is not complete brittle fracture. Although these values should be further verified using finite element analysis, it to some extent explains the reason why the nucleation site of cleavage cracks appears at ~0.6 mm away from the ductile crack tips (Fig. 7a). The stretch zone that is characterized by quite a number of fine dimples (e.g., Fig. 6c) is considered as the most effective zones to consume the impact energy. Thus, some empirical correlations between the stretch zone size and crack initiation toughness have been built [11,27,33]. Here, a simple correlation in Eq. (5) is used to explain which method is more reasonable to determine dynamic fracture toughness (JId). SZW = 89(JId /E) (5) where SZW is the stretch zone width. For as-rolled specimens, the SZW is ~2 mm (Fig. 6a). The calculated JId is ~4719 kJ/m2, even much higher than the JId obtained by the maximum load point method. This means that the stretch zone contains the ductile propagation stage after the maximum load. While for the CGHAZ specimens, the calculated JId is ~429 kJ/m2, which lies in between 8
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Fig. 8. (a and b) SEM micrographs showing the morphology of cleavage fracture for the ICCGHAZ specimen, (c and d) optical micrographs showing the propagation of main cleavage cracks and the initiation sites of secondary cleavage cracks, as arrowed.
the values (see Fig. 4) obtained by the compliance changing rate method [8] and averaging method of yield and peak loads [9]. And the latter method gives an acceptable conservative estimation of JId. As a result, this method is most suitable to directly determine dynamic fracture toughness. From this work, the ICCGHAZ remarkably deteriorates the dynamic fracture toughness because numerous MA constituents formed during the second intercritically welding thermal cycle assist the crack initiation. Thus, to protect the investigated steel weldment from low stress brittle fracture in the structural application, suitable post-welding heat treatments, e.g., high temperature tempering or annealing, should be employed to eliminate these micro-constituents. 4. Conclusions This work investigated the dynamic fracture toughness of three different heat-treated specimens using instrumental Charpy impact tester. The crack initiation point was determined with four different methods based on impact load-deflection curves. It can be concluded that the JId value is overestimated when the maximum load is regarded as the crack initiation point, especially in the ductile fracture mode; while it is over-conservative when the initiation point is determined at the time of ti/tmax = 0.4. The JId value decreases in the order of as-rolled state > CGHAZ > ICCGHAZ, irrespective of the estimation methods. Based on the present data, the ratio of Chary impact energy to JId3 drops gradually with weld thermal cycles (that is, the ratio is 0.35 for as-rolled specimens, 0.18 for CGHAZ, and 0.09 for ICCGHAZ). The initiation of ductile macrocrack occurs prior to the maximum load; whereas, for the ICCGHAZ the micro-cleavage cracks may initiate from the blocky MA constituents. The SZW confirmed that the estimated JId value is most acceptable when the crack initiation point is determined at the point of (Py + Pm)/2. Acknowledgement This work is supported by the National Natural Science Foundation of China (No. 51605084 and U1708265) and the Fundamental Research Funds for the Central Universities of China (N170304019 and N170308028). 9
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[23] Puppala G, Moitra A, Sathyanarayanan Kaul SR, Sasikala G. Evaluation of fracture toughness and impact toughness of laser rapid manufactured Inconel-625 structures and their co-relation. Mater Des 2014;59:509–15. [24] Lambert-Perlade A, Gourgues AF, Besson J, Sturel T, Pineau A. Mechanism and modeling of cleavage fracture in simulated heat-affected zone microstructures of a high-strength low alloy steel. Metall Mater Trans A 2004;35A:1039–53. [25] Lan LY, Yu M, Qiu CL. On the local mechanical properties of isothermally transformed bainite in low carbon steel. Mater Sci Eng A 2019;742:442–50. [26] Mohseni P, Solberg JK, Karlsen M, Akselsen OM, Østby E. Investigation of mechanism of cleavage fracture initiation in intercritically coarse grained heat affected zone of HSLA steel. Mater Sci Technol 2012;28:1261–8. [27] Sreenivasan PR, Ray SK, Vaidyanathan S, Rodriguez P. Measurement of stretch zone height and its relationship to crack tip opening displacement and initiation J-value in an AISI 316 stainless steel. Fatigue Fract Eng Mater Struct 1996;19:855–68. [28] Cao R, Li J, Liu DS, Ma JY, Chen JH. Micromechanism of decrease of impact toughness in coarse-grain heat-affected zone of HSLA steel with increasing welding heat input. Metall Mater Trans A 2015;46A:2999–3014. [29] Li XD, Ma XP, Subramanian SV, Shang CJ. EBSD characterization of secondary microcracks in the heat affected zone of a X100 pipeline steel weld joint. Int J Fract 2015;193:131–9. [30] Fairchild DP, Howden DG, Clark WAT. The mechanism of brittle fracture in a microalloyed steel: Part I. Inclusion- induced cleavage. Metall Mater Trans A 2000;31A:641–52. [31] Lan LY, Qiu CL, Song HY, Zhao DW. Correlation of martensite-austenite constituent and cleavage crack initiation in welding heat affected zone of low carbon bainitic steel. Mater Lett 2014;125:86–8. [32] Davis CL, King JE. Cleavage initiation in the intercritically reheated coarse-grained heat-affected zone: Part I. fractographic evidence. Metall Mater Trans A 1994;25A:563–73. [33] Tagawa T, Kayamori Y, Hira H. Statistical scatter of fracture toughness in the ductile-brittle transition of a structural steel. Eng Fract Mech 2010;77:3077–86.
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Engineering Fracture Mechanics journal homepage: www.elsevier.com/locate/engfracmech
Dynamic fracture behavior of low carbon bainitic steel after different welding thermal cycles
T
⁎
Liangyun Lana,b, , Yiting Zhanga, Xiangwei Konga,b a
School of Mechanical Engineering and Automation, Northeastern University, Shenyang 110819, China Key Laboratory of Vibration and Control of Aero-Propulsion Systems, Ministry of Education of China, Northeastern University, Shenyang 110819, China
b
A R T IC LE I N F O
ABS TRA CT
Keywords: Dynamic fracture toughness Low carbon steel Stretch zone width Crack initiation point
Dynamic fracture toughness is an important parameter to assess the fracture resistance at crack initiation although the actual crack initiation point is hard to be precisely detected. In this work, three different heat treatment microstructures (i.e., hot-rolled, welding thermal simulation coarse grained HAZ (CGHAZ) and intercritically reheated CGHAZ microstructures) were studied by instrumented Charpy impact tester. The dynamic fracture toughness was obtained based on impact load-deflection curves using various estimation methods. The dynamic fracture toughness decreased with the applied welding thermal cycles irrespective of the estimation methods. The stretch zone width proved that the fracture toughness is most acceptable when the crack initiation point is considered as the middle point between the general yield load and maximum load in the load-deflection curves.
1. Introduction Welding is one of the most common methods in industrial practice for joining components because of low costs and high efficiency. However, it also yields many defects, e.g., heterogeneous microstructure, local brittle zone, cold cracks and residual stress. These defects threat the integrity and reliability of weldment structure [1]. Thus, the attention from metallurgists has been always paid to the reasons that cause the deterioration in mechanical properties of welded joints. Neves and Loureiro [2] showed that a decrease in toughness was observed in the weld metal and heat-affected zone (HAZ) with increasing heat input as a result of the increasing size of the local brittle zone sampled by fatigue crack front. Yang et al. [3] found that the fracture toughness of fusion zone is poorest in the X80 pipeline weldment and the hard and brittle martensite-austenite (MA) constituents may play roles of the fracture origins. Linear elastic fracture mechanics based plane strain fracture toughness (KIC) approach or Elastic-plastic fracture mechanics based J-integral approach is recognized as an applicable method to carry out structural integrity assessment [1–3]. Determinations of these static fracture toughness parameters, however, require very strict and somewhat complex procedures [4] such as large and thick samples, sharp fatigue pre-crack with reasonable initial length. Particularly for the HAZ with size limits and irregular shape, the sample preparation is hard to meet these requirements. Charpy impact test is also widely applied to reveal the impact fracture behaviors of materials in industry due to the simple, rapid and inexpensive testing procedure. However, from the fracture mechanics point of view, the impact absorbed energy itself cannot be directly used for quantitative safety assessment [5]. Because impact specimen only has a Charpy V or U notch, the total impact energy should be composed of the crack initiation and
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Corresponding author at: School of Mechanical Engineering and Automation, Northeastern University, Shenyang 110819, China. E-mail addresses: [email protected] (L. Lan), [email protected] (X. Kong).
https://doi.org/10.1016/j.engfracmech.2019.106653 Received 7 April 2019; Received in revised form 29 August 2019; Accepted 30 August 2019 Available online 30 August 2019 0013-7944/ © 2019 Elsevier Ltd. All rights reserved.
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propagation energies and the ratio of them mainly depends on the stress-strain state that the samples undergo. Besides the mathematical relationships established between the fracture toughness and Charpy impact energy [6,7], dynamic fracture toughness that is derived from the impact load-deflection curve is another important parameter to assess the fracture resistance at crack initiation [8,9]. However, the crack initiation point is hard to be detected exactly even using instrumented impact tester. In general, the peak load that occurs in the load-deflection curve is approximately regarded as the onset point of crack propagation [10]. It is noted that this method intentionally improves the crack initiation energy, especially under the condition of ductile fracture. Several other methods were proposed in literature to ascertain the crack initiation point, such as compliance changing rate method [8] and averaging method of yield and peak loads [9]. As a result, the dynamic fracture toughness (JId) can be estimated with these methods. Zhu et al. [11] considered that the JId is acceptable when the crack initiation point is considered to occur at a load of (Pm + Py)/2, where Pm and Py is maximum and general yield load, respectively. Moitra et al. [12] showed that the JId follows a power law relationship with the inter-particle spacing for different regions in the HAZ of a nuclear grade steel. Although the dynamic fracture toughness reflects the resistance of materials to macro-crack initiation at the impact load condition, majority of studies only focused on the effect of microstructure on Charpy impact energy rather than dynamic fracture toughness [13–16]. As we known, the Charpy impact energy cannot be directly used for the structural integrity assessment. Meanwhile, the total absorbed energy does not characterize the resistance of materials to crack initiation or propagation at a certain stage because different stages of fracture have some different physical phenomena [17]. In this work, three different heat-treated microstructures, i.e., as-rolled state, welding thermal simulated-CGHAZ (coarse grained HAZ) and ICCGHAZ (intercritically reheated CGHAZ), were employed to study the relationship between microstructure and dynamic fracture toughness. In combination with the morphology of fracture surface, the underlying mechanism that deteriorates the dynamic fracture toughness of the ICCGHAZ was explored.
2. Experimental methods The chemical composition of studied low carbon steel was 0.053 C, 0.25 Si, 1.68 Mn, 0.45 (Cr + Mo), 0.086 (Nb + V + Ti), 0.53 (Cu + Ni), 0.0012 B (Wt.%) balanced by Fe, similar to the composite of X80 pipeline steel [3]. Hot rolled steel plates with a thickness of 13 mm were produced using thermo-mechanically controlled processing (TMCP) in our laboratory and the detailed TMCP parameters as well as as-rolled mechanical properties were reported in Ref. [18]. Here, some important curves were reused to further confirm dynamic fracture parameters and these values were compared with as-welded dynamic fracture parameters. Specimens with a dimension of 11 × 11 × 55 mm were cut from the hot rolled steel plates with their longest dimension perpendicular to rolling direction. Single and double-pass welding thermal cycles were simulated to obtain CGHAZ and ICCGHAZ microstructures, respectively, employing a resistance heating thermo-simulator. For the first welding thermal cycle, the peak temperature was set at 1350 °C to make the austenite grain coarsening; while for the second cycle, the peak temperature was set at 740 °C that is just higher than the α to γ phase transformation temperature (i.e., Ac1 ~ 710 °C). The cooling rate for both cycles was adopted as a cooling time (t8/5 is the elapsed time when the temperature dropped from 800 °C to 500 °C) of 30 s, which is supposed to simulate submerged arc welding with the heat input of about 38 kJ/mm. Fig. 1 shows the measured temperature curve for double-pass welding thermal cycle. Each of simulation processes was conducted at least three times using the fresh specimens. After the welding simulation, all specimens were re-machined into the standard size (10 × 10 × 55 mm) of Charpy impact samples and a 2-mm deep V-notch was machined at the region of interest. The impact testing was performed at −20 °C in an instrumented drop weight impact tester equipped with an oscilloscope. The actual load–deflection curves were recorded with the oscilloscope. The microstructures as well as the morphology of second cracks underneath the fracture surface were examined by a Leica DMIRM optical microscopy after the metallographic specimens were prepared by conventional methods, i.e. manual and mechanical polishing, 3% nital etching and hot air drying. Microhardness measurement was conducted using a FM700 hardnesstesting machine, employed a 0.49 N load. At least 7 indenters were detected for each kind of heat treated microstructures and average
Fig. 1. Measured temperature profile for simulating double-pass welding thermal cycles (P1 and P2: two peak temperatures). 2
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Fig. 2. Raw data impact load- deflection curves for the specimens with different heat treatment states (a) as-rolled state, (b) CGHAZ, (c) ICCGHAZ (three curves given in each condition represent three independent results).
values were reported below. A Zeiss Ultra 55 field emission scanning electron microscope (SEM) was used to observe the morphology of fracture surfaces of broken Charpy specimens.
3. Results and discussion 3.1. Impact load-deflection curves Fig. 2 represents raw impact load-deflection curves of the specimens after subjected to different heat treatments. These curves can be sub-divided into five different stages, i.e., elastic and plastic deformation, ductile propagation stage, brittle (unstable) fracture stage, and ductile fracture stage [10,19], although not all of these stages exist on each curve. For the hot-rolled specimens, their curves always exhibited a relatively stable fracture stage, implying that the ductile crack propagation occurs [20]. Occasionally, some of them also have a very short stage of unstable crack extension (as arrowed in Fig. 2a). In contrast, all of the welding thermally treated specimens exhibited predominantly brittle unstable fracture behaviors since the loads abruptly dropped to zero immediately after the peak load (Fig. 2b and c). This behavior not only leads to a huge reduction in sample’s deflection but also makes the phenomenon of load fluctuations more obvious. Furthermore, the double-pass welding thermal cycle (Fig. 2c) seems to decrease the peak loads as compared to Fig. 2b. To determine the dynamic fracture parameters accurately, these oscillations should be eliminated through data smoothing.
3.2. Dynamic fracture toughness Fig. 3a shows an example of the raw data curve and corresponding smoothened curve. Some dynamic fracture parameters such as Py, Pm and corresponding absorbed energy can be directly obtained from the smoothing curve, as marked with symbols. Fig. 3b represents one group of typical smooth curves for different heat treated specimens. The general yield stress (σy) for full size CVN specimens can be estimated from the Py using the following equation [8]: 3
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Fig. 3. (a) Comparison of raw curve and smoothing curve and the determination of dynamic fracture parameter, (b) one typical group of smooth curves for different heat treated specimens.
σy =
2.99·Py·W B·(W − a0)2
(1)
where B is the sample thickness, W is the sample width and a0 is the initial crack length. For standard Charpy V notch specimens, the B = W = 10 mm, a0/W = 0.2. Based on elastic-plastic fracture mechanics, the JId can be evaluated from the energy absorbed up to the crack initiation point as [8,12],
JId =
ηEa B (W − a0)
(2)
Here, Ea is the absorbed energy up to the crack initiation point and η = 1.45 for standard Charpy V notch specimens. Therefore, the most important procedure to obtain the JId is how to determine the crack initiation point. For simplicity, the crack initiation point is assumed to be at the maximum load point even with an obvious plastic deformation (i.e., Ea = Ei) [21]. Kobayashi [8] developed the compliance changing rate method to determine the crack initiation point and found that the ratio of initiation to maximum load energies is about 0.8 (i.e., Ea/Ei = 0.8). This result is verified by Tronskar et al. [22] using a laser interferometer to measure displacement. Ghoneim and Hammad [9] proposed that the crack initiation is more reasonably taken as at the load point of (Py + Pm)/2 because the crack initiation often occurs above the yield load but prior to maximum load. Zhu et al. [11] and Puppala et al. [23] proved that the JId integral estimated by this method largely agrees with the experimental results. Norris [24] reported the crack initiates at the time of ti/tmax = 0.4, where ti is the time to crack initiation; tmax is the time to reach the maximum load, based on a comparison of finite element results with experimental data. Moitra et al. [12] confirmed that this method gives more conservative estimation of JId than the averaging method of yield and maximum loads. When the Py is almost equal to Pm (e.g. the ICCGHAZ curve in Fig. 3b), this means that the testing temperature approaches to the nil ductility temperature of the specimen. Thus, the fracture stress (σf) can be approximately evaluated from the σy using the following expression:
σf = Cf σy
(3)
where the stress intensification factor Cf is taken as 2.57 for the full size CVN specimens [21]. The details of dynamic fracture parameters determined from curves are shown in Table 1 and the JId values calculated with various methods are presented in Fig. 4, where JId1 represents that the crack initiation occurs at maximum load point, JId2 crack initiation at the potential energy 0.8 times Ei, JId3 crack initiation at the load point of (Py + Pm)/2, and JId4 crack initiation at the time of ti/tmax = 0.4. It can be seen that for these different methods the crack initiation points determined are not the same, which leads to the fact that the dynamic fracture toughness is quite different, especially for the hot rolled specimen. Obviously, the maximum load point represents an over-estimated JId value because it may contain some stable ductile crack propagations after plastic deformation, which Table 1 Average values of dynamic fracture parameters for all impact specimens. Specimen
Py (kN)
Pm (kN)
Ei (J)
Ep (J)
σy (MPa)
σf (MPa)
Hot rolled CGHAZ ICCGHAZ
17.8 17.6 15.6
22.8 22.3 16.8
65 39 8
85 11 2.9
831 822 696
– 2112 1788
4
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Fig. 4. The calculated dynamic fracture toughness with various methods.
will be verified with SEM observation later. The crack initiation point at the time of ti/tmax = 0.4 gives the most conservative JId values as compared to other methods. However, for the ICCGHAZ, the crack initiation point determined with this method is even much lower than the general yield point. It seems to be unreasonable since the crack initiation is hard to occur at so lower stress condition. Compared with the fracture toughness (256–438 kJ/m2 at different regions for X80 pipeline welded joint) obtained by Yang et al.[3], the JId3 value is most acceptable to characterize the resistance of different heat-treated specimens to cracking. In spite of the estimation methods used, the JId always decreases in order of Hot-rolled, CGHAZ, and then ICCGHAZ. This means that the weld thermal cycle deteriorates the fracture toughness. Meanwhile, it can be found that the ratio of Charpy impact energy to JId gradually decreases due to the effect of weld thermal cycles. That is, using the JId3 values the ratio is 0.35 for hot-rolled specimens, 0.18 for CGHAZ, and 0.09 for ICCGHAZ. These phenomena are in good agreement with the results by Lambert-Perlade et al. [24], who considered that the microstructural evolution under the welding thermal cycles is responsible for the deterioration of fracture toughness. 3.3. Microstructures The microstructures with different heat treatments are shown in Fig. 5. The hot rolled specimens have a mixed degenerate upper bainite and lath bainite. The pancake-shape prior austenite grain boundaries can be observed clearly (arrowed in Fig. 5a) and their thickness measured by line intercept method is ~10.4 um. After the specimens were subject to the welding thermal cycles, these pancaked prior austenite grain boundaries disappear and they become near polygonal-shape (arrowed in Fig. 5b) with a coarse average grain size of ~86 um, although the main microstructure is still transformed into upper bainite (or bainitic ferrite). Each coarse prior austenite grain seems to be partitioned into several bainitic packets because the packet morphology is described by the paralleling distribution of carbides and stringer-like MA constituents inside each a bainitic packet [25]. When the specimens were further subjected to the intercritically reheated thermal cycles, the matrix microstructure does not show a notable change except for slightly tempering effect (Fig. 5c) as compared to Fig. 5b. However, it is evident that the prior austenite grain boundaries are decorated by numerous black particles that are considered to be massive MA constituents (as arrowed) and their maximum width ranges from ~3 to 5.6 μm measured at different views of field. These islands are directly derived from the newly formed austenite during the second welding thermal cycles because the peak temperature is just slightly higher than Ac1. And they inherently possess much higher hardness and brittleness than the ferritic matrix due to higher amount of carbon atoms enriched on them [24–26]. Although some difference exists in microstructural morphology for these three different heat treated specimens, the microhardness does not exhibit an obvious change. The average values in HV are 263 ± 12, 270 ± 10, and 268 ± 18 for the hot rolled, CGHAZ and ICCGHAZ specimens, respectively. 3.4. SEM fractography and its relation to fracture toughness The morphology of fracture surface not only reveals the micro-mechanism of fracture but also verifies the mechanical properties by measuring some characteristic parameters, e.g., stretch zone width and unit cleavage facet size [11,24,27–29]. Fig. 6a gives a typical example of macrograph for the fracture surface of as-rolled specimen. The lateral expansion event can be found and the fracture surface is macroscopically rough with the presence of unstable fracture region in the center of the surface (as cycled region). The fibrous stretch zone with a width of ~2 mm appears between the V-notch and unstable fracture region, as marked with double arrows. These two regions correspond to two different fracture mechanisms, i.e., ductile and cleavage fracture shown in Fig. 6c and b, respectively. In combination with Fig. 2a, the change of fracture mechanism from ductile fracture to cleavage fracture is mainly responsible for a short stage of fast decreasing load (arrowed) because the cleavage fracture belongs to a typical low stress brittle fracture [20]. Meanwhile, it also implies that the ductile propagation stage corresponds to the fibrous stretch zone prior to unstable 5
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Fig. 5. Microstructure at different heat treatment states (a) hot-rolled state (RD: rolling direction), (b) CGHAZ, (c) ICCGHAZ.
fracture region. However, the initiation point of ductile macrocrack still cannot be fixed before or after the peak load. In addition, these fine cleavage facets with elongated shape along the rolling direction (Fig. 6b) reflect the crystallographic orientation distribution of microstructure because only high-angle boundaries can effectively arrest the propagation of cleavage cracks [15,24]. Fig. 7 shows the fracture surface of the CGHAZ specimen. From the macrograph (Fig. 7a), the fracture surface is rather flat compared to Fig. 6a and the phenomenon of lateral expansion does not exist. Based on macro-fracture patterns the entire flat fracture seems to initiate at the location near to the notch, as marked with a black grid. The magnified morphology of this square region is shown in Fig. 7b. It is interesting to find that the real initiation site is derived from a coarse TiN precipitate with the size of about 2.7 um (shown in the inset) that is located inside a prior austenite grain. The dotted lines and arrows indicate the propagation paths and directions of cleavage fracture. This precipitate is likely to be cracked by itself to stimulate the initiation of microcrack as the tessellated residual stresses between the TiN and the matrix increase the cleavage-initiating potency of TiN inclusions [30,31]. More interestingly, a very narrow ligament (the width ~0.182 mm) can be found at the bottom of Fig. 7a and it is separated from the cleavage fracture region by a straight line (arrowed in Fig. 7b). This means that the CGHAZ specimen first fractured with ductile tearing for a very short while, and then the cleavage initiation was triggered because the stress concentration region enlarges with the growth of ductile crack [5]. Here, the cleavage initiation site is ~0.8 mm away from the blunt notch. As the cleavage cracks propagate very fast, the front of ductile cracks encounters the cleavage cracks at the same time, which forms an obvious boundary between them. In combination with Fig. 2b, it can be inferred that the ductile tearing must occur to before the peak load because no ductile propagation stage appears after the peak load. That is to say, the increasing load suddenly dropped to zero immediately after the occurrence of cleavage fracture. Under a higher magnification, the unit cleavage facets can be largely identified according to the traces of ridges as delineated in Fig. 7c because these ridges formed are to compensate the local misorientation between two neighboring cleavage facets [17]. The reinitiation site for each facet is likely located at the interior of the grains based on the river-pattern, and the size of unit cleavage facets seems to be comparable to the prior austenite grain size of the CGHAZ (Fig. 5b). For the ICCGHAZ specimen, the morphology of fracture surface is characterized by coarse cleavage facets without obvious ductile tearing region (Fig. 8a). Meanwhile, it can be seen that most of coarse cleavage facets are re-initiated attached to prior austenite grain boundaries as the ridges of cleavage facet are converged at the grain boundary, as arrowed in Fig. 8b. This implies that the microstructural feature at the prior austenite grain boundaries for the ICCGHAZ is in favor of assisting the initiation of microcracks. 6
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Fig. 6. Fracture surface of as-rolled specimens: (a) macrograph, (b) morphology of cleavage fracture at the center of macro- fracture surface and RD indicates rolling direction, (c) morphology of ductile fracture near to the V-notch position.
Davis and King [32] showed that the brittle fracture initiation at prior austenite grain boundaries is attributed to the fact that debonding of these necklace-type MA constituents with matrix microstructure is aggravated under the overlap of residual phase transformation stresses and the stress concentrations. Mohseni et al. [26] confirmed that the initiation sites of cleavage facets are characterized by twinned martensite and argued that the initiation site is hard to be identified due to the divergence of river pattern in the lath-like microstructure. Here, it is acceptable that some particles located at the initiation site of cleavage facets (arrowed in Fig. 8a) are considered to be the MA constituents since they have a similar morphology, size and location to the MA constituents appeared in Fig. 5c. The morphology of secondary cleavage cracks underneath the fracture surface can further illustrate the initiation sites of cleavage crack because the stress state that they undergo may be similar to the main fracture surface. Fig. 8c shows a portion of cross-sectional microstructure in the vicinity of the fracture surface. The arrows in Fig. 8c represent the unit crack paths during main crack propagation. The unit path is defined as the region in which the crack propagates almost in a straight line [29]. It is closely related to the size of cleavage facets. Evidently, most prior austenite grain boundaries inhibit the linear propagation of main cracks effectively although the cracks, sometimes, exhibit some deviations inside the prior austenite grains. Many secondary cracks can be found in the same or nearby prior austenite grains of the main cracks, as marked with a square grid in Fig. 8c. This squared region is magnified in Fig. 8d to show the detailed morphology of secondary cracks. They are approximately parallel to the main fracture surface and most of them seem to initiate individually at massive MA constituents along the prior austenite grain boundaries, as marked with arrows. Especially for these small microcracks, they seem just nucleate on the MA constituents and do not further propagate probably due to fast decreasing stress. Thus, these blocky MA constituents along the prior austenite grain boundaries indeed act as potential crack initiation sites in the ICCGHAZ. Because the ICCGHAZ sample almost exhibits complete brittle fracture, the fracture stress calculated from the general yield load 7
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Fig. 7. Fracture surface of the CGHAZ specimens: (a) macrograph, (b) the initiation of cleavage fracture at the TiN precipitate near to the V notch position, (c) coarse cleavage facets.
(Table 1) should be satisfied with the formation of critical Griffith microcrack size according to the classic Griffith theory: 1/2 π Eγp ⎞ σc = ⎜⎛ ⎟ 2 ⎝ (1 − ν ) d ⎠
(4)
where E, ν, and γp are the material’s parameters, i.e., Young’s modulus, Poisson’s ratio and the effective surface energy of microcrack, respectively. Using the same values given in Ref. [31], the critical Griffith microcrack size is ~ 3.2 μm at the maximum fracture stress. This critical size agrees well with the width size of MA constiutents formed in the ICCGHAZ mentioned above. In contrast, coarse precipitates such as TiN in the CGHAZ, no doubt, acts as a trigger for cleavage fracture (Fig. 7b). It can be calculated that the critical fracture stress needed is ~1938 MPa as the TiN size measured on the SEM image is ~2.7 μm. This required critical stress is lower than the calculated fracture stress (Table 1) in the front of ductile cracks. It is acceptable because the CGHAZ is not complete brittle fracture. Although these values should be further verified using finite element analysis, it to some extent explains the reason why the nucleation site of cleavage cracks appears at ~0.6 mm away from the ductile crack tips (Fig. 7a). The stretch zone that is characterized by quite a number of fine dimples (e.g., Fig. 6c) is considered as the most effective zones to consume the impact energy. Thus, some empirical correlations between the stretch zone size and crack initiation toughness have been built [11,27,33]. Here, a simple correlation in Eq. (5) is used to explain which method is more reasonable to determine dynamic fracture toughness (JId). SZW = 89(JId /E) (5) where SZW is the stretch zone width. For as-rolled specimens, the SZW is ~2 mm (Fig. 6a). The calculated JId is ~4719 kJ/m2, even much higher than the JId obtained by the maximum load point method. This means that the stretch zone contains the ductile propagation stage after the maximum load. While for the CGHAZ specimens, the calculated JId is ~429 kJ/m2, which lies in between 8
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Fig. 8. (a and b) SEM micrographs showing the morphology of cleavage fracture for the ICCGHAZ specimen, (c and d) optical micrographs showing the propagation of main cleavage cracks and the initiation sites of secondary cleavage cracks, as arrowed.
the values (see Fig. 4) obtained by the compliance changing rate method [8] and averaging method of yield and peak loads [9]. And the latter method gives an acceptable conservative estimation of JId. As a result, this method is most suitable to directly determine dynamic fracture toughness. From this work, the ICCGHAZ remarkably deteriorates the dynamic fracture toughness because numerous MA constituents formed during the second intercritically welding thermal cycle assist the crack initiation. Thus, to protect the investigated steel weldment from low stress brittle fracture in the structural application, suitable post-welding heat treatments, e.g., high temperature tempering or annealing, should be employed to eliminate these micro-constituents. 4. Conclusions This work investigated the dynamic fracture toughness of three different heat-treated specimens using instrumental Charpy impact tester. The crack initiation point was determined with four different methods based on impact load-deflection curves. It can be concluded that the JId value is overestimated when the maximum load is regarded as the crack initiation point, especially in the ductile fracture mode; while it is over-conservative when the initiation point is determined at the time of ti/tmax = 0.4. The JId value decreases in the order of as-rolled state > CGHAZ > ICCGHAZ, irrespective of the estimation methods. Based on the present data, the ratio of Chary impact energy to JId3 drops gradually with weld thermal cycles (that is, the ratio is 0.35 for as-rolled specimens, 0.18 for CGHAZ, and 0.09 for ICCGHAZ). The initiation of ductile macrocrack occurs prior to the maximum load; whereas, for the ICCGHAZ the micro-cleavage cracks may initiate from the blocky MA constituents. The SZW confirmed that the estimated JId value is most acceptable when the crack initiation point is determined at the point of (Py + Pm)/2. Acknowledgement This work is supported by the National Natural Science Foundation of China (No. 51605084 and U1708265) and the Fundamental Research Funds for the Central Universities of China (N170304019 and N170308028). 9
Engineering Fracture Mechanics 220 (2019) 106653
L. Lan, et al.
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