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Effect of Ni on the hot ductility and hot cracking susceptibility of high Mn austenitic cast steel
Materials Science & Engineering A 618 (2014) 295–304
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Effect of Ni on the hot ductility and hot cracking susceptibility of high Mn austenitic cast steel Kyutae Han a, Jaehong Yoo a, Bongkeun Lee b, Ilwook Han b, Changhee Lee a,n a b
Department of Materials Science and Engineering, Hanyang University, Seoul 133-791, Republic of Korea POSCO Technical Research Laboratories, 1, Pohang 790-785, Republic of Korea
art ic l e i nf o
a b s t r a c t
Article history: Received 30 June 2014 Received in revised form 7 September 2014 Accepted 8 September 2014 Available online 16 September 2014
In the present study, the hot ductility behavior and hot cracking susceptibility of high manganese cast steel were investigated. Based on the 18Mn–0.6C alloy, a nickel was added for cryogenic impact toughness by maintaining the low temperature austenite stability to prevent strain induced martensitic transformation considering the appropriate stacking fault energy for mechanical twinning. A hot ductility test and Varestraint test were carried out to clarify the behavior of heat affected zone (HAZ) liquation and ductility dip cracking that were observed at the multi-pass weld heat affected zone of a nickel added high manganese weld metal for cryogenic uses. The heating and cooling rates were calculated by the three-dimensional Rosenthal's equation based on the actual welding heat input. A brittle temperature range, and a critical strain temperature range for each alloy showed no significant difference, but the overall ductility of the nickel added alloy was lower than the 18Mn alloy due to the lower degree of dynamic recrystallization by the higher stacking fault energy and the existence of the M3P/γ eutectic and MnS formed by the severe P segregation and higher S content, respectively. As a result of the longitudinal Varestraint test, the total and maximum HAZ crack length and cracked HAZ length of the nickel added alloy were larger than those of 18Mn alloy. The solid-state fracture, i.e. the ductility dip cracking was observed both in the on-cooling hot ductility tested and Varestraint tested alloys. & 2014 Elsevier B.V. All rights reserved.
Keywords: Cryogenic high manganese alloy Liquation crack Ductility dip crack Hot ductility test Varestraint test
1. Introduction As conventional energy sources are steadily exhausted, demands for the development of new and clean energy sources have increased. A recent report has indicated that neutral gas is the fastest growing major fuel with global demand rising by 60% from 2010 to 2040 [1]. Therefore, it has been generally accepted that the development of suitable cryogenic alloys for liquefied natural gas (LNG) tanks is one of the most important matter with increasing demand for the transportation of LNG. Austenitic high manganese steel is a strong candidate for replacing conventional cryogenic materials such as 9Ni steels, austenitic stainless steels, aluminum and nickel alloys and Invar steels [2]. Choi et al. [3] also reported that the cryogenic application of high manganese steel can be considered for the construction of LNG tanks due to its competitiveness in price and low temperature properties. Most cryogenic alloys have a fully face centered cubic (FCC) crystal structure that does not transit from ductile to brittle fracture even at low temperatures. Therefore, the high manganese
n
Corresponding author. Tel.: þ 82 2 2220 0388; fax: þ 82 2 2299 0389. E-mail address: [email protected] (C. Lee).
http://dx.doi.org/10.1016/j.msea.2014.09.040 0921-5093/& 2014 Elsevier B.V. All rights reserved.
alloys for cryogenic uses should have a fully austenitic microstructure with the low temperature toughness requirements. In the view of the development of high manganese welding consumables or cast alloys, inhomogeneous weld metal and cast microstructures by manganese evaporation, alloying element segregation and depletion during welding and solidification can severely affect the degradation of mechanical properties of the final product [4]. Especially, the formation of ε-martensite due to the depletion of austenitic stabilizing elements, i.e., carbon, manganese, degrades cryogenic mechanical properties compare to the case of the fully austenitic microstructure [4–6]. Meanwhile, in the view of the weld metal hot cracking susceptibility, the low carbon content is generally better than the higher carbon content considering all of possible cracks, i.e., a solidification crack, a liquation crack and a ductility dip crack in the fusion zone and the heat affected zone in welds. However, the carbon content extensively affects the austenite stability, i.e., the stacking fault energy [7–9], therefore it should be maintained as an appropriate content for fully austenitic microstructure of high manganese steel. Accordingly, the addition of alloying elements that can increase the stacking fault energy, for example, nickel [10] and aluminum [11–13] should be considered, instead of the carbon addition for cryogenic austenitic stability and cracking susceptibility of high
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manganese weld metals and cast alloys. As a result of Ni addition to 18Mn–0.6C based weld metal, the impact toughness at 196 1C was remarkably increased from 28 J to 47 J of the 18 wt%Mn– 0.6 wt%C, less than 18 wt%Mn- more than 5 wt%Ni–0.6 wt%C containing weld metals, respectively by the fully mechanical twinning instead of the stain induced martensitic transformation (SIMT). Recently, a number of researchers [14–21] reported cryogenic high manganese steels and welding consumables with respect to the mechanical properties, deformation and fracture mechanisms but the research for the hot cracking susceptibility has rarely done. Sutton et al. [22] reported the solidification cracking susceptibility by the cast pin tear test of welding consumables with a variation of manganese, carbon and aluminum contents. However, in view of the alloy design and optimization for the development of the welding consumables, liquation crack and ductility dip crack susceptibility can be more practical factor than the solidification cracking susceptibility considering the weld heat affected zone crack which is more frequently observed in the heat affected zone by subsequent passes in the multi-pass actual welds. In this study, the effect of nickel on the hot ductility behaviors and hot cracking susceptibility of cryogenic high manganese cast alloys were investigated for clarifying the mechanism of the micro crack formation that was observed in the weld heat affected zone of nickel added multi pass welded specimen considering the eutectic formation temperature and degree of dynamic recrystallization by a hot ductility test and Varestraint test. In addition, the 18 wt%Mn–0.6 wt%C alloy which was investigated with commercial austenitic stainless steels in the view of hot cracking susceptibility in the previous study [23] was selected for comparing the hot crack susceptibility with the nickel added alloy. Various hot cracking criteria and detailed microstructural analysis via a hot ductility test and Varestraint test were carried out for the exact evaluation of crack susceptibility of both fully austenitic alloys, the nickel added high manganese alloy and the Ni-free 18Mn– 0.6C alloy.
2. Experimental procedures 2.1. Materials characterization of the as-cast material Two cast alloys, a 18Mn-cast (18 wt%Mn–0.6 wt%C) and Niadded cast (o18 wt% Mn, 4 5 wt% Ni, 0.6 wt% C) were manufactured via a vacuum melting process for a hot ductility test and a Varestraint test. The chemical compositions of the two cast specimens are shown in Table 1. The distribution of phases and constituents were analyzed using an optical microscope, a scanning electron microscope (SEM), and an energy dispersive X-ray spectroscope (EDS). The extracted specimens were fine polished with diamond powder with a particle size of 1 μm. Chemical etching was conducted using 98 vol% ethanol and 2 vol% nitric acid.
were extracted only in the center of the cast ingot for equiaxed and randomly oriented grains to prevent preferential cracking due to the different grain orientations with respect to the tensile axis and a different chemical composition if specimens were extracted from different areas. A heat-affected zone (HAZ) simulation of the weld HAZ and high temperature tensile tests were carried out using a Gleeble thermo-mechanical simulator. The free span (distance between the two copper jigs) was selected as 10 mm to minimize the temperature fluctuation that can occur during resistance heating. K-type (Chromel-alumel) thermo-couples were directly attached to specimens via percussion welding. A thermal cycle for simulating a weld coarse grained heat affected zone (CGHAZ) was calculated using a three dimensional Rosenthal's equation [24] using directly measured thermal conductivity (λ), 17.18 W/mK and thermal diffusivity (a), 6.38 10 6 m2/s by a C-THERM TCI™ thermal conductivity analyzer. The heating and cooling rates were 300 1C/s and 50 1C/s, respectively, as a result of the calculation. A series of hot ductility tests were carried out in the temperature range of 800–1350 1C. In the case of the onheating test, the strain was loaded with a strain rate of 50 mm/sec at each test temperature. On-cooling tests were performed from the peak temperature of the zero ductility temperature (ZDT), which was determined at the on-heating test. After exposure to the peak temperature of ZDT, the strain was loaded at each test temperature during cooling. Hot ductility tests were carried out in an argon atmosphere for preventing oxidation of the fractured surface. 2.3. Varestraint test An additional hot cracking test was carried out to evaluate the susceptibility of the actual weld heat affected zone cracking using a Varestraint method [25]. The test employs the rectangular specimens with the dimensions of 127 25.4 3 mm3. The autogeneous gas tungsten arc welding (GTAW) specification for the Varestraint test is shown in Table 2. Welding conditions for the GTAW were chosen to a weld pool penetration depth that was half the thickness of the specimen. A 4% augmented strain was applied perpendicular to the welding direction immediately after the torch passed (a longitudinal mode). All of specimens for the Varestraint test were extracted from the same area where hot ductility specimens were taken from.
3. Results and discussion 3.1. Microstructure of as-cast materials
The test employs a cylindrical specimen, 3 mm in diameter and 90 mm long with 10 mm threads on both ends. All of the samples
Fig. 1 shows the microstructures and results of detailed analysis of precipitates and eutectic constituents of as-cast 18Mn and Niadded alloys used in this study. The cellular dendritic size of the Ni-added alloy was slightly smaller than the 18Mn as shown in Fig. 1(a) and (d). The solidification temperature range of both alloys was nearly the same as shown in the Fe–C quasi binary diagram (Fig. 2) calculated by the Thermo-Calc. Therefore, it could be concluded that the difference of precipitates and its distribution determined the grain and sub grain size of each specimen. In the Ni-added alloy, randomly-distributed coarse globular or rod-type
Table 1 Chemical composition of the specimens (wt%).
Table 2 The welding specification for a Varestraint test.
2.2. Hot ductility test
Specimen
C
Si
Mn
Ni
P
S
Ti
Fe
18Mn Ni-added
0.59 0.61
Added Added
18.5 o 18
0.0 5o
0.02 0.02
0.02 0.04
0.0 0.014
Bal. Bal.
Process Current (A)
Voltage (V)
Travel speed (mm/s)
Arc length (mm)
Shielding gas/flow rate (Lmm 1)
GTAW
12
4
2
Argon, 20
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Fig. 1. As-cast microstructure and inclusions: (a), (b) 18Mn-cast, (d)–(f) Ni-cast, (c(a)) and (c(b)) EDS of M3C/γ and M3P/γ eutectic, respectively.
Fig. 2. Fe–C quasi-binary diagrams: (a) 18Mn-cast, (b) Ni-cast alloys.
manganese sulfides exist at inner grain and grain boundaries due to the higher S content. In addition, intragranular titanium carbide also exists by the small titanium content as shown in Fig. 1(e). Therefore, the two precipitates, numerous and randomly distributed MnS and TiC that precipitated above a solidus temperature near 1300 1C, reduced the dendritic cell size of the Ni-added alloy (118 μm for Ni-added alloy, 130 μm for 18Mn-alloy, respectively) by the precipitation pinning effect on dendritic growth and small and higher concentration of γ dendrite nuclei. The carbon content that can precipitate as a M3C carbide after titanium carbide precipitation was 0.575 and 0.59 wt% in the Ni-added-cast and 18Mn-cast, respectively, as a result of the stoichiometric calculation considering carbon arrest by titanium in the Ni-cast specimen (Table 3). Consequently, nearly the same quantities of the M3C and M3C/γ eutectics could exist in both cast alloys (Fig. 1(b), 1(c (a)), 1(e)). Additionally, numerous M3P/γ eutectic constituents were observed only in the Ni-added alloy even though the phosphorus
Table 3 The calculation of carbon contents available for M3C, M3C/γ formation considering an arrestment of carbon by Ti. wt%
C
Ti
C–Ti/4
18Mn Ni-added
0.59 0.61
0 0.014
0.59 0.575
contents of the 18Mn and Ni-added-cast alloys were nearly the same as 0.02 wt% (Fig. 1(c (b)), 1(e), 1(f)). It is well known that the solubility of phosphorus in austenite is small. Therefore, phosphorus can be easily segregated at grain boundaries or sub grain boundaries, resulting in the formation of Fe3P inclusion or the γ þ Fe3P eutectic if the phosphorus contents reaches as high as 1 wt% at 1050 1C [26]. Additionally, with the presence of nickel in austenitic alloys, the solubility of phosphorus and sulfur is
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diminished due to the practically zero solubility in pure nickel [27]. Kaneko et al. also reported that the solubility of phosphorus in the fully austenitic alloy is decreased by Ni addition throughout the thermodynamic study of Fe–P–C and Fe–P–Ni systems [28]. In the case of the Ni-added alloy used in this study, more than 5 wt% nickel was added for low temperature austenite stability, resulting in the prevention of strain induced ε-martensite formation and hence higher impact toughness at the cryogenic condition with less than 18 wt% manganese and 0.6 wt% carbon. It could be concluded that the additional nickel in fully austenitic high alloyed manganese steels can significantly lower the phosphorus solubility in the matrix compared to nickel-free austenitic high manganese steel, resulting in more severe phosphorus segregation in the grain or sub-grain boundaries and M3P/γ eutectic formation during solidification. Severe phosphorus segregation was due to the larger dendritic cell and grain size of the as-cast specimen compared to the as-welded case. However, in the view of the reducing phosphorus-related hot crack, the phosphorus content should be sufficiently limited in the Ni-added fully austenitic high manganese steel.
3.2. Results of the hot ductility test The hot ductility behavior during on-heating and on-cooling cycles of 18Mn and Ni-added alloys is shown in Fig. 3(a) and (b), respectively. The overall ductility of the Ni-added at each test temperature was less than the 18Mn in both on-heating and oncooling cases. The zero ductility temperature (ZDT) of both specimens was the same as 1300 1C, which is near the solidus temperatures of the specimens (Fig. 2). The brittle temperature range (BTR), the temperature difference between ZDT and oncooling ductility recovery starting temperature, is also same as 200 1C. Additionally, the critical strain temperature (CST) was also measured for both alloys. CST is the temperature at critical strain that is smaller than restrained, thermally induced strain in which fracture can occur [29]. Authors recommended that the CST can be determined as the temperature at a 7% reduction of area during the on-cooling test if the strain cannot be directly measured through their several experiments. Above the CST, i.e. 1070 1C and 1075 1C for the 18Mn and the Ni-added alloy, respectively, the material fracture strain is less than the thermal strain during welding, resulting in higher susceptibility to cracking. The critical strain temperature range (CSTR), the temperature difference between ZDT and CST, which is the other criteria of hot crack susceptibility regarding the restraint strain for both alloys [29],
is nearly similar at approximately 230 1C. Even though several criteria for hot cracking susceptibility of both specimens are similar, it should be analyzed that the discrepancy in the overall hot ductility, and a huge ductility drop at a comparatively low temperature in the Ni-added alloy, and a ductility drop at an intermediate temperature during cooling for both specimens.
3.2.1. The hot ductility at intermediate temperatures Figs. 4–8 show the microstructure and fractured surface of the hot ductility tested 18Mn and Ni-added at each test temperature, respectively. In the temperature range of 800–900 1C during heating, the ductility dropped as shown in the on-heating ductility curves (the solid square line of Fig. 3(a) and (b)). The hot ductility of a material is governed by the balance between work hardening and dynamic softening. If work hardening predominates, intergranular cracking occurs as a result of grain boundary sliding that promotes void nucleation and coalescence at grain boundaries. As the temperature increased from 800 to 900 1C, the resistance of the grain interior to plastic flow decreased following grain boundary sliding, which leads to early void formation [30]. The degree of ductility drop in the Ni-added alloy at the temperature range of 800–900 1C was more severe than it was in the 18Mn (Fig. 3) due to the large amount of the γ/M3P eutectic constituents at the grain boundaries and coarse MnS inclusions (Fig. 1(e), (f)). Zhang et al. [26] also reported that the thermo plasticity of high manganese austenitic steel can be reduced by the eutectic constituent and MnS inclusions. If the coarse and irregular shape of MnS and eutectic constituents exist on the grain or sub grain boundaries, the inclusion/matrix interface can be the initiation source of a crack due to the stress concentration that accelerates the fracture. As shown in Figs. 4(a) and 5(a), dynamic recrystallized grains were observed near fractured area of both specimens of the on-heating tested at 900 1C. A material that has a low stacking fault energy like an austenitic stainless steel, such as a high manganese austenitic steel, more likely undergoes dynamic recrystallization instead of dynamic recovery as a softening process during hot deformation [29,30]. Figs. 4(a) and 5(a) show round shaped cracks that were arrested and isolated from the initial grain boundaries by the dynamic recrystallized grains. As shown in Figs. 4(d) and 5(d), fractured surfaces consisted of flat surfaces with shallow dimples due to grain boundary sliding at early stage of the crack formation and dimples due to ductile transgranular failure. The dynamic recrystallized grain nucleates in highly -strained grain boundaries, isolating and blunting the initial crack, resulting in improved hot ductility. As shown in the hot
Fig. 3. Hot ductility curves, brittle temperature range (BTR)s, zero ductility temperature (ZDT)s, critical strain temperature (CST)s, and critical strain temperature range (CSTR)s: (a) 18Mn-cast, (b) Ni-cast alloys.
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Fig. 4. The microstructure and fractured surface of the 18Mn-cast specimens on-heating hot ductility tested at: (a) and (d) 900 1C, (b) and (e) 1200 1C, (c) and (f) 1300 1C.
Fig. 5. The microstructure and fractured surface of the Ni-cast specimens on-heating hot ductility tested at: (a) and (d) 900 1C, (b) and (e) 950 1C, (c) and (f) 1200 1C, respectively.
ductility curves (Fig. 3), the ductility increased in the 18Mn alloy and the change of the ductility curvature was observed in the Niadded alloy after initiation of the dynamic recrystallization at 900 1C. Even though the onset temperature of dynamic recrystallization was 900 1C for both alloys, the degree of recrystallization was different. As shown in Figs. 4(a) and 5(a), there were more partially recrystallized grains in the 18Mn than the Ni-added alloy. Dynamic recrystallization occurs when the applied strain is greater than the critical strain. The Zener–Hollomon parameter (Z) which
describes the high temperature strain of a material: Z ¼ έ exp Q =RT
ð1Þ
where έ is the strain rate, Q is the activation energy for deformation, R is the gas constant, and T is the absolute temperature. As shown in Eq. (1), the critical strain for dynamic recrystallization depends on the deformation temperature, the strain rate and chemical composition dependent activation energy [31]. Therefore, the effect of grain size difference [30] and fine titanium
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Fig. 6. (a) M3C/γ eutectic at vicinity of the crack in Fig. 4(b) (open circle), (b) M3P/γ eutectic at vicinity of the crack in Fig. 5(b) (open circle), and (c) M3C/γ eutectic at vicinity of the crack in Fig. 5(c) (open circle).
Fig. 7. The microstructure and fractured surface of the on-cooling hot ductility tested 18Mn-cast specimens: (a) and (d) microstructure and fractured surface of 1200 1C, (b) microstructure of 1000 1C, (c) and (e) microstructure and fractured surface of 900 1C.
carbide particles (o0.5 μm) [32] of the Ni-added alloy can influence the onset temperature of the dynamic recrystallization and retard the growth of the nuclei of the dynamic recrystallization, are limited due to the slight difference in grain size and small quantities of TiC precipitates. Consequently, for the condition of the same test temperature and strain rate, the chemical composition of each material is the primary factor for the dynamic recrystallization. In view of the chemical composition difference between 18Mn and Ni-added alloys, the manganese and nickel contents varied as 18.5 wt% Mn and 0 wt% Ni in the 18Mn-cast specimen and less than 18 wt%Mn, and more than 5 wt%Ni in the Ni-cast specimen while the carbon content is same as 0.6 wt%. Cabaňas et al. [31] and Hamada et al. [33] reported that manganese delays dynamic recrystallization in the Fe–Mn binary alloys and high manganese TWIP steel due to the increase in the activation energy for hot working with increasing manganese content. In addition, solutes that increase the stacking fault energy (SFE) will favor dynamic recovery, while solutes that reduce the stacking fault energy will favor dynamic recrystallization as the primary softening mechanism. As a result of the thermo-dynamic calculation considering alloying elements and temperature using thermodynamic databases and equations [34–36], the stacking fault
energy of the Ni-added was greater than the 18Mn alloy with a difference of 4–5 mJ/m2 throughout the wide temperature range. Though no exact data was available for the high temperature stacking fault energy, it was expected to increase with temperature and the SFE difference between the 18Mn and the Ni-added alloy would be maintained. Consequently, the degree of dynamic recrystallization in the Ni-added alloy was smaller than the nickelfree 18Mn alloy due to the increase in stacking fault energy resulting in retarded dynamic recrystallization by the nickel addition. The flow stress level of two cast materials can also be explained by Mn–C dipoles short range ordering due to the high carbon content, as studied by Saeed-Akbari et al. [37]. The theoretical ordering index (TOI), which was defined for providing a measurement of the fraction of carbon and manganese atoms in different Fe–Mn–C systems [37], was 0.15 and 0.17 for the 18Mn and the Ni-added alloys, respectively. For TOI values between 0.1 and 0.3, the population of CMn6 octahedral clusters that can act as obstacles for dislocations increased. Therefore, as shown in Figs. 4(a), 7(b) and Figs. 5(a), 8(b), the dynamic recrystallization of both specimens were limited and only partial dynamic recrystallized grains were observed due to the high flow stress level of the two cast materials at a fast strain rate condition.
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Fig. 8. The microstructure and fractured surface of the on-cooling hot ductility tested Ni-cast specimens: (a) microstructure of 1200 1C, (b) microstructure of 1050 1C, (c) and (d) microstructure and fractured surface of 900 1C, respectively.
3.2.2. The liquation cracking As the temperature increased below the eutectic temperature of γ/M3C, the hot ductility of the 18Mn increased slightly, but the Ni-added alloy gradually decreased in the temperature range of 900–1000 1C even though dynamic recrystallization occurred after 900 1C. As shown in Figs. 5(b), 6(b), and 5(e), the melting trace was observed along grain boundaries and the fractured surface due to the (Fe, Mn)3P/γ eutectic melting occurred at 950 1C. Fig. 6(b) is the M3P/γ eutectic, with an enlarged image of the open circle shown in Fig. 5(b). Mintz et al. reported that the M3P/γ eutectic temperature can be lowered to 950 1C with the highly concentrated carbon at the interdendritic area and grain boundaries. Consequently, the ductility drop at 950 1C of the Ni-added alloy was due to the melting of the M3P/γ eutectic constituent at the grain boundaries that were formed in the as-cast condition. Consequently, in view of the hot ductility behavior, the nickel addition to the high manganese austenitic steel seems to have a detrimental effect of retarding dynamic recrystallization and the low melting eutectic constituent, M3P/γ formation by reducing the solubility of phosphorus in the austenite matrix. As the temperature increases to the M3C/γ eutectic temperature of the each specimen (Fig. 3), the ductility abruptly decreases by eutectic melting of M3C/γ, resulting in the formation of the liquid film on the grain boundaries that accelerates the change in fracture mode from transgranular to intergranular as shown in Figs. 4(b), 6(a), 4 (e), and Figs. 5(c), 6(c), 5(f). Fig. 6(a) and (c) are M3C/γ eutectics, enlarged images of open circles in Figs. 4(b) and 5(c), respectively. As the temperature approaches the zero ductility temperature (ZDT), the degree of ductility drop of both specimens gently decreases while the liquid penetrated and wetted all over the grain boundaries as shown in Fig. 4(c) and (f) at 1300 1C (ZDT). At this stage, no dynamic recrystallization occurred because the
fracture strain was lower than the critical strain for the initiation of recrystallization. The ductility of this stage is a function of the extent of liquid formation via incipient melting at Si, P and S segregated grain boundaries and liquid continuity along the grain boundaries [29].
3.2.3. The on-cooling behavior On-cooling ductility recovery of both alloys started under 1100 1C near the eutectic temperature of γ/M3C. Figs. 7(a), (d), and 8(a) are typical microstructures near the fractured region and fracture surface of 18Mn and Ni-added alloys tested on-cooling at 1200 1C, respectively. The microstructures of both specimens were similar to those in the ductility gently decreased stage of the onheating test. The difference is the wider grain boundaries due to the prolonged exposure near the ZDT, resulting in extensive liquid formation along the grain boundaries. Even though most of the melted grain boundaries have started to solidify at this temperature range under the equilibrium solidus temperature of both alloys (Fig. 2), but as shown in the fractography at 1200 1C in Fig. 7(d), boundaries were penetrated by the M3C/γ eutectic liquid in a wetted condition at the moment of fracture. With a further decrease in temperature below the eutectic temperature of M3C/γ, the ductility increased due to the depletion of the liquid along the grain boundaries with the prevention of the initial crack by recrystallized grains and grain boundary migration (Figs. 3, 8(a), and (b), microstructures of the on-cooling tested 18Mn at 1000 1C and Ni-added alloy at 1050 1C, respectively.). As shown in Fig. 3, a ductility drop was observed in the temperature range of 1000 1C to 900 1C for the 18Mn and 1050 1C to 900 1C for the Ni-added alloy, respectively. When compared to microstructures of the on-heating tested specimens, preciously melted and resolidified “ghost
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boundaries [29]” were clearly observed as shown in Figs. 7(c) and 8 (c). Authors also mentioned that the previously melted and resolidified boundaries were enriched in solute elements, including P, S and Si, compared to newly formed boundaries resulting in significant embrittlement, which are favored for crack initiation and propagation along the boundaries. No recrystallized grains and relatively long and straight cracks along the previously melted and resolidified grains were observed in the on-cooling both alloys at 900 1C (Figs. 7(c), 8(c)). The fractographies also show no liquidrelated evidence, but the completely intergranular solid-state fracture along resolidified boundaries with some proof, i.e. a striation or ripped eutectics on the fracture (Figs. 7(e) and 8(d)). This ductility dip phenomena can be influenced by the peak temperature of the on-cooling test in terms of the eutectic melting, quantities of the liquid formation, degree of grain boundary liquid penetration, the grain boundary migration and the final grain size that can affect the dynamic recrystallization. With the on-cooling temperature of ZDT, no dynamic recrystallization or ductility dip phenomena of either cast alloy was observed near at 900 1C as seen in the multi-pass weld heat affected zone of the Ni added aswelded specimen in this study. If the peak temperature increases to the zero strength temperature (ZST), the recrystallization temperature can be shifted to a higher temperature, and vice versa. This is primarily due to the effect of the grain size on dynamic recrystallization and the quantities of eutectic constituents and possible precipitation that can affect the embrittlement or hardening of the matrix.
3.3. Results of the Varestraint test Fig. 9 shows the results of the Varestraint tests for the 18Mn and the Ni-added alloys. As shown in Figs. 10 and 11, all of cracks were counted and measured at the HAZ of the autogeneous gas tungsten arc welded 18Mn and Ni-added alloy. The total crack length (TCL) and maximum crack length (MCL) those were measured in direction perpendicular to the fusion line of Niadded alloy were slightly larger than those of the 18Mn, and the TCL and MCL of 18Mn alloys were nearly the same as the TCL and MCL of 18Mn–0.6C weld metal HAZ with the 4% augmented strain of the previous study [23]. Besides the conventional TCL and MCL measurements, the other criterion designed by Lee [38], termed “cracked HAZ length (CHL)”, measured along the fusion line in the HAZ where cracks formed was also measured to directly compare the hot ductility test evaluating HAZ cracking susceptibility along the same peak temperature region parallel to the fusion line. The CHL, the length parallel to the fusion line between the start point of the cooling portion of the weld pool and the region where the
Fig. 10. Cracked HAZ length (CHL) of 18Mn-cast and Ni-cast Varestraint tested specimens.
last crack was observed, of the 18Mn and the Ni-added alloy were 3.6 and 5.5 mm, respectively (Figs. 10, 11). Based on the directly-measured, on-cooling hot crack susceptibility test for the actual welds, the CHL results can be another good indicator for the evaluation of specimens with similar conventional criterions. Fig. 8 shows microstructures of the Varestraint tested specimens. Isothermal lines from the fusion line are indicated in Fig. 11(a) and (d) as the calculation results of the three dimensional Rosenthal's equation. As shown in these figures, cracks in the HAZ were initiated from the fusion line and propagated towards the base metal until the lowest eutectic temperature of each specimen, 1080 1C of M3C/γ and 950 1C of M3P/γ for the 18Mn and Ni-added alloys, respectively. Fig. 11(b) and (e) are the enlarged images of the M3C/γ eutectic for the 18Mn and Ni-added alloys, while Fig. 11(f) is the M3P/γ eutectics near the end tip of the crack in the Ni-added alloy. The ductility dip crack along the migrated grain boundaries were also observed below the lowest eutectic temperature as shown in Fig. 11(c) and (d). Consequently, cracks initiated above the solidus or eutectic temperature by a liquation mechanism and propagated along the liquid-penetrated grain boundaries followed the solid-state crack formation due to the resolidified and embrittled boundaries with the same crack formation mechanism of the as-welded alloys.
4. Conclusions The hot cracking susceptibility of two as-cast alloys with different manganese and nickel contents were investigated by the hot ductility and the Varestraint tests for clarifying the mechanism of cracks in a multi-pass weld heat affected zone of the cryogenic high manganese steel. Conclusions of this study are as follows.
Fig. 9. Results of the Varestraint test: total crack length, maximum crack length, cracked HAZ length.
1. The additional nickel on fully austenitic high alloyed manganese steels can significantly reduce the phosphorus solubility in the matrix compared to nickel-free austenitic high manganese steel, resulting in more severe phosphorus segregation in the grain or sub grain boundaries and M3P/γ eutectic formation during solidification in the as-cast condition.
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Fig. 11. Microstructure of the Varestraint tested specimens: (a) low magnification image of the fusion zone and heat affected zone of 18Mn-cast specimen, (b) enlarged image of the open square area in Fig. (a), (c) crack across the line of Te along migrated grain boundary of 18Mn-cast specimen weld HAZ, (d) low magnification image of the fusion zone and heat affected zone of the Ni-cast specimen, (e) and (f) enlarged image of the solid and open square area in Fig. (d), respectively.
2. Hot cracking susceptibility criteria from the hot ductility test, BTR and CSTR of 18Mn and Ni-added cast alloys were similar but the overall ductility of the Ni-added alloy was less than that of the 18Mn alloy due to the low temperature of eutectic melting of the γ/M3P, coarse and randomly distributed MnS inclusions, and the low degree of the dynamic recrystallization by the higher stacking fault energy compared to the 18Mn alloy. 3. Dynamic recrystallization was limited due to the high carbon contents of both 18Mn and Ni-added alloys resulting in only partial dynamic recrystallization under a condition of fast strain rate. 4. Significant ductility drops were observed in both 18Mn and Ni-added alloys as a result of M3C/γ eutectic melting at each eutectic temperatures during on-heating tests. That was attributed to a change of fracture mode from transgranular to intergranular and was accelerated by the liquid films formation along grain boundaries. 5. With the on-cooling peak temperature of ZDT, no dynamic recrystallization following the ductility dip phenomena of both cast alloys were observed near 900 1C as observed in the multipass weld heat affected zone of the Ni added welds. 6. Total and maximum crack lengths and the cracked HAZ length that allow direct measurements of the on-cooling hot crack susceptibility of the actual welds for the Ni-added alloy was bigger than the 18Mn alloy as a result of the Varestraint test.
Acknowledgment The author wishes to thank the POSCO Technical Research Laboratory for supporting this research (2012Z026). References [1] ExxonMobil, The Outlook for Energy: A View to 2040, Texas, Exxon Mobile Corporation, 2012, pp. 1–57.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Ni on the hot ductility and hot cracking susceptibility of high Mn austenitic cast steel Kyutae Han a, Jaehong Yoo a, Bongkeun Lee b, Ilwook Han b, Changhee Lee a,n a b
Department of Materials Science and Engineering, Hanyang University, Seoul 133-791, Republic of Korea POSCO Technical Research Laboratories, 1, Pohang 790-785, Republic of Korea
art ic l e i nf o
a b s t r a c t
Article history: Received 30 June 2014 Received in revised form 7 September 2014 Accepted 8 September 2014 Available online 16 September 2014
In the present study, the hot ductility behavior and hot cracking susceptibility of high manganese cast steel were investigated. Based on the 18Mn–0.6C alloy, a nickel was added for cryogenic impact toughness by maintaining the low temperature austenite stability to prevent strain induced martensitic transformation considering the appropriate stacking fault energy for mechanical twinning. A hot ductility test and Varestraint test were carried out to clarify the behavior of heat affected zone (HAZ) liquation and ductility dip cracking that were observed at the multi-pass weld heat affected zone of a nickel added high manganese weld metal for cryogenic uses. The heating and cooling rates were calculated by the three-dimensional Rosenthal's equation based on the actual welding heat input. A brittle temperature range, and a critical strain temperature range for each alloy showed no significant difference, but the overall ductility of the nickel added alloy was lower than the 18Mn alloy due to the lower degree of dynamic recrystallization by the higher stacking fault energy and the existence of the M3P/γ eutectic and MnS formed by the severe P segregation and higher S content, respectively. As a result of the longitudinal Varestraint test, the total and maximum HAZ crack length and cracked HAZ length of the nickel added alloy were larger than those of 18Mn alloy. The solid-state fracture, i.e. the ductility dip cracking was observed both in the on-cooling hot ductility tested and Varestraint tested alloys. & 2014 Elsevier B.V. All rights reserved.
Keywords: Cryogenic high manganese alloy Liquation crack Ductility dip crack Hot ductility test Varestraint test
1. Introduction As conventional energy sources are steadily exhausted, demands for the development of new and clean energy sources have increased. A recent report has indicated that neutral gas is the fastest growing major fuel with global demand rising by 60% from 2010 to 2040 [1]. Therefore, it has been generally accepted that the development of suitable cryogenic alloys for liquefied natural gas (LNG) tanks is one of the most important matter with increasing demand for the transportation of LNG. Austenitic high manganese steel is a strong candidate for replacing conventional cryogenic materials such as 9Ni steels, austenitic stainless steels, aluminum and nickel alloys and Invar steels [2]. Choi et al. [3] also reported that the cryogenic application of high manganese steel can be considered for the construction of LNG tanks due to its competitiveness in price and low temperature properties. Most cryogenic alloys have a fully face centered cubic (FCC) crystal structure that does not transit from ductile to brittle fracture even at low temperatures. Therefore, the high manganese
n
Corresponding author. Tel.: þ 82 2 2220 0388; fax: þ 82 2 2299 0389. E-mail address: [email protected] (C. Lee).
http://dx.doi.org/10.1016/j.msea.2014.09.040 0921-5093/& 2014 Elsevier B.V. All rights reserved.
alloys for cryogenic uses should have a fully austenitic microstructure with the low temperature toughness requirements. In the view of the development of high manganese welding consumables or cast alloys, inhomogeneous weld metal and cast microstructures by manganese evaporation, alloying element segregation and depletion during welding and solidification can severely affect the degradation of mechanical properties of the final product [4]. Especially, the formation of ε-martensite due to the depletion of austenitic stabilizing elements, i.e., carbon, manganese, degrades cryogenic mechanical properties compare to the case of the fully austenitic microstructure [4–6]. Meanwhile, in the view of the weld metal hot cracking susceptibility, the low carbon content is generally better than the higher carbon content considering all of possible cracks, i.e., a solidification crack, a liquation crack and a ductility dip crack in the fusion zone and the heat affected zone in welds. However, the carbon content extensively affects the austenite stability, i.e., the stacking fault energy [7–9], therefore it should be maintained as an appropriate content for fully austenitic microstructure of high manganese steel. Accordingly, the addition of alloying elements that can increase the stacking fault energy, for example, nickel [10] and aluminum [11–13] should be considered, instead of the carbon addition for cryogenic austenitic stability and cracking susceptibility of high
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manganese weld metals and cast alloys. As a result of Ni addition to 18Mn–0.6C based weld metal, the impact toughness at 196 1C was remarkably increased from 28 J to 47 J of the 18 wt%Mn– 0.6 wt%C, less than 18 wt%Mn- more than 5 wt%Ni–0.6 wt%C containing weld metals, respectively by the fully mechanical twinning instead of the stain induced martensitic transformation (SIMT). Recently, a number of researchers [14–21] reported cryogenic high manganese steels and welding consumables with respect to the mechanical properties, deformation and fracture mechanisms but the research for the hot cracking susceptibility has rarely done. Sutton et al. [22] reported the solidification cracking susceptibility by the cast pin tear test of welding consumables with a variation of manganese, carbon and aluminum contents. However, in view of the alloy design and optimization for the development of the welding consumables, liquation crack and ductility dip crack susceptibility can be more practical factor than the solidification cracking susceptibility considering the weld heat affected zone crack which is more frequently observed in the heat affected zone by subsequent passes in the multi-pass actual welds. In this study, the effect of nickel on the hot ductility behaviors and hot cracking susceptibility of cryogenic high manganese cast alloys were investigated for clarifying the mechanism of the micro crack formation that was observed in the weld heat affected zone of nickel added multi pass welded specimen considering the eutectic formation temperature and degree of dynamic recrystallization by a hot ductility test and Varestraint test. In addition, the 18 wt%Mn–0.6 wt%C alloy which was investigated with commercial austenitic stainless steels in the view of hot cracking susceptibility in the previous study [23] was selected for comparing the hot crack susceptibility with the nickel added alloy. Various hot cracking criteria and detailed microstructural analysis via a hot ductility test and Varestraint test were carried out for the exact evaluation of crack susceptibility of both fully austenitic alloys, the nickel added high manganese alloy and the Ni-free 18Mn– 0.6C alloy.
2. Experimental procedures 2.1. Materials characterization of the as-cast material Two cast alloys, a 18Mn-cast (18 wt%Mn–0.6 wt%C) and Niadded cast (o18 wt% Mn, 4 5 wt% Ni, 0.6 wt% C) were manufactured via a vacuum melting process for a hot ductility test and a Varestraint test. The chemical compositions of the two cast specimens are shown in Table 1. The distribution of phases and constituents were analyzed using an optical microscope, a scanning electron microscope (SEM), and an energy dispersive X-ray spectroscope (EDS). The extracted specimens were fine polished with diamond powder with a particle size of 1 μm. Chemical etching was conducted using 98 vol% ethanol and 2 vol% nitric acid.
were extracted only in the center of the cast ingot for equiaxed and randomly oriented grains to prevent preferential cracking due to the different grain orientations with respect to the tensile axis and a different chemical composition if specimens were extracted from different areas. A heat-affected zone (HAZ) simulation of the weld HAZ and high temperature tensile tests were carried out using a Gleeble thermo-mechanical simulator. The free span (distance between the two copper jigs) was selected as 10 mm to minimize the temperature fluctuation that can occur during resistance heating. K-type (Chromel-alumel) thermo-couples were directly attached to specimens via percussion welding. A thermal cycle for simulating a weld coarse grained heat affected zone (CGHAZ) was calculated using a three dimensional Rosenthal's equation [24] using directly measured thermal conductivity (λ), 17.18 W/mK and thermal diffusivity (a), 6.38 10 6 m2/s by a C-THERM TCI™ thermal conductivity analyzer. The heating and cooling rates were 300 1C/s and 50 1C/s, respectively, as a result of the calculation. A series of hot ductility tests were carried out in the temperature range of 800–1350 1C. In the case of the onheating test, the strain was loaded with a strain rate of 50 mm/sec at each test temperature. On-cooling tests were performed from the peak temperature of the zero ductility temperature (ZDT), which was determined at the on-heating test. After exposure to the peak temperature of ZDT, the strain was loaded at each test temperature during cooling. Hot ductility tests were carried out in an argon atmosphere for preventing oxidation of the fractured surface. 2.3. Varestraint test An additional hot cracking test was carried out to evaluate the susceptibility of the actual weld heat affected zone cracking using a Varestraint method [25]. The test employs the rectangular specimens with the dimensions of 127 25.4 3 mm3. The autogeneous gas tungsten arc welding (GTAW) specification for the Varestraint test is shown in Table 2. Welding conditions for the GTAW were chosen to a weld pool penetration depth that was half the thickness of the specimen. A 4% augmented strain was applied perpendicular to the welding direction immediately after the torch passed (a longitudinal mode). All of specimens for the Varestraint test were extracted from the same area where hot ductility specimens were taken from.
3. Results and discussion 3.1. Microstructure of as-cast materials
The test employs a cylindrical specimen, 3 mm in diameter and 90 mm long with 10 mm threads on both ends. All of the samples
Fig. 1 shows the microstructures and results of detailed analysis of precipitates and eutectic constituents of as-cast 18Mn and Niadded alloys used in this study. The cellular dendritic size of the Ni-added alloy was slightly smaller than the 18Mn as shown in Fig. 1(a) and (d). The solidification temperature range of both alloys was nearly the same as shown in the Fe–C quasi binary diagram (Fig. 2) calculated by the Thermo-Calc. Therefore, it could be concluded that the difference of precipitates and its distribution determined the grain and sub grain size of each specimen. In the Ni-added alloy, randomly-distributed coarse globular or rod-type
Table 1 Chemical composition of the specimens (wt%).
Table 2 The welding specification for a Varestraint test.
2.2. Hot ductility test
Specimen
C
Si
Mn
Ni
P
S
Ti
Fe
18Mn Ni-added
0.59 0.61
Added Added
18.5 o 18
0.0 5o
0.02 0.02
0.02 0.04
0.0 0.014
Bal. Bal.
Process Current (A)
Voltage (V)
Travel speed (mm/s)
Arc length (mm)
Shielding gas/flow rate (Lmm 1)
GTAW
12
4
2
Argon, 20
100
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Fig. 1. As-cast microstructure and inclusions: (a), (b) 18Mn-cast, (d)–(f) Ni-cast, (c(a)) and (c(b)) EDS of M3C/γ and M3P/γ eutectic, respectively.
Fig. 2. Fe–C quasi-binary diagrams: (a) 18Mn-cast, (b) Ni-cast alloys.
manganese sulfides exist at inner grain and grain boundaries due to the higher S content. In addition, intragranular titanium carbide also exists by the small titanium content as shown in Fig. 1(e). Therefore, the two precipitates, numerous and randomly distributed MnS and TiC that precipitated above a solidus temperature near 1300 1C, reduced the dendritic cell size of the Ni-added alloy (118 μm for Ni-added alloy, 130 μm for 18Mn-alloy, respectively) by the precipitation pinning effect on dendritic growth and small and higher concentration of γ dendrite nuclei. The carbon content that can precipitate as a M3C carbide after titanium carbide precipitation was 0.575 and 0.59 wt% in the Ni-added-cast and 18Mn-cast, respectively, as a result of the stoichiometric calculation considering carbon arrest by titanium in the Ni-cast specimen (Table 3). Consequently, nearly the same quantities of the M3C and M3C/γ eutectics could exist in both cast alloys (Fig. 1(b), 1(c (a)), 1(e)). Additionally, numerous M3P/γ eutectic constituents were observed only in the Ni-added alloy even though the phosphorus
Table 3 The calculation of carbon contents available for M3C, M3C/γ formation considering an arrestment of carbon by Ti. wt%
C
Ti
C–Ti/4
18Mn Ni-added
0.59 0.61
0 0.014
0.59 0.575
contents of the 18Mn and Ni-added-cast alloys were nearly the same as 0.02 wt% (Fig. 1(c (b)), 1(e), 1(f)). It is well known that the solubility of phosphorus in austenite is small. Therefore, phosphorus can be easily segregated at grain boundaries or sub grain boundaries, resulting in the formation of Fe3P inclusion or the γ þ Fe3P eutectic if the phosphorus contents reaches as high as 1 wt% at 1050 1C [26]. Additionally, with the presence of nickel in austenitic alloys, the solubility of phosphorus and sulfur is
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diminished due to the practically zero solubility in pure nickel [27]. Kaneko et al. also reported that the solubility of phosphorus in the fully austenitic alloy is decreased by Ni addition throughout the thermodynamic study of Fe–P–C and Fe–P–Ni systems [28]. In the case of the Ni-added alloy used in this study, more than 5 wt% nickel was added for low temperature austenite stability, resulting in the prevention of strain induced ε-martensite formation and hence higher impact toughness at the cryogenic condition with less than 18 wt% manganese and 0.6 wt% carbon. It could be concluded that the additional nickel in fully austenitic high alloyed manganese steels can significantly lower the phosphorus solubility in the matrix compared to nickel-free austenitic high manganese steel, resulting in more severe phosphorus segregation in the grain or sub-grain boundaries and M3P/γ eutectic formation during solidification. Severe phosphorus segregation was due to the larger dendritic cell and grain size of the as-cast specimen compared to the as-welded case. However, in the view of the reducing phosphorus-related hot crack, the phosphorus content should be sufficiently limited in the Ni-added fully austenitic high manganese steel.
3.2. Results of the hot ductility test The hot ductility behavior during on-heating and on-cooling cycles of 18Mn and Ni-added alloys is shown in Fig. 3(a) and (b), respectively. The overall ductility of the Ni-added at each test temperature was less than the 18Mn in both on-heating and oncooling cases. The zero ductility temperature (ZDT) of both specimens was the same as 1300 1C, which is near the solidus temperatures of the specimens (Fig. 2). The brittle temperature range (BTR), the temperature difference between ZDT and oncooling ductility recovery starting temperature, is also same as 200 1C. Additionally, the critical strain temperature (CST) was also measured for both alloys. CST is the temperature at critical strain that is smaller than restrained, thermally induced strain in which fracture can occur [29]. Authors recommended that the CST can be determined as the temperature at a 7% reduction of area during the on-cooling test if the strain cannot be directly measured through their several experiments. Above the CST, i.e. 1070 1C and 1075 1C for the 18Mn and the Ni-added alloy, respectively, the material fracture strain is less than the thermal strain during welding, resulting in higher susceptibility to cracking. The critical strain temperature range (CSTR), the temperature difference between ZDT and CST, which is the other criteria of hot crack susceptibility regarding the restraint strain for both alloys [29],
is nearly similar at approximately 230 1C. Even though several criteria for hot cracking susceptibility of both specimens are similar, it should be analyzed that the discrepancy in the overall hot ductility, and a huge ductility drop at a comparatively low temperature in the Ni-added alloy, and a ductility drop at an intermediate temperature during cooling for both specimens.
3.2.1. The hot ductility at intermediate temperatures Figs. 4–8 show the microstructure and fractured surface of the hot ductility tested 18Mn and Ni-added at each test temperature, respectively. In the temperature range of 800–900 1C during heating, the ductility dropped as shown in the on-heating ductility curves (the solid square line of Fig. 3(a) and (b)). The hot ductility of a material is governed by the balance between work hardening and dynamic softening. If work hardening predominates, intergranular cracking occurs as a result of grain boundary sliding that promotes void nucleation and coalescence at grain boundaries. As the temperature increased from 800 to 900 1C, the resistance of the grain interior to plastic flow decreased following grain boundary sliding, which leads to early void formation [30]. The degree of ductility drop in the Ni-added alloy at the temperature range of 800–900 1C was more severe than it was in the 18Mn (Fig. 3) due to the large amount of the γ/M3P eutectic constituents at the grain boundaries and coarse MnS inclusions (Fig. 1(e), (f)). Zhang et al. [26] also reported that the thermo plasticity of high manganese austenitic steel can be reduced by the eutectic constituent and MnS inclusions. If the coarse and irregular shape of MnS and eutectic constituents exist on the grain or sub grain boundaries, the inclusion/matrix interface can be the initiation source of a crack due to the stress concentration that accelerates the fracture. As shown in Figs. 4(a) and 5(a), dynamic recrystallized grains were observed near fractured area of both specimens of the on-heating tested at 900 1C. A material that has a low stacking fault energy like an austenitic stainless steel, such as a high manganese austenitic steel, more likely undergoes dynamic recrystallization instead of dynamic recovery as a softening process during hot deformation [29,30]. Figs. 4(a) and 5(a) show round shaped cracks that were arrested and isolated from the initial grain boundaries by the dynamic recrystallized grains. As shown in Figs. 4(d) and 5(d), fractured surfaces consisted of flat surfaces with shallow dimples due to grain boundary sliding at early stage of the crack formation and dimples due to ductile transgranular failure. The dynamic recrystallized grain nucleates in highly -strained grain boundaries, isolating and blunting the initial crack, resulting in improved hot ductility. As shown in the hot
Fig. 3. Hot ductility curves, brittle temperature range (BTR)s, zero ductility temperature (ZDT)s, critical strain temperature (CST)s, and critical strain temperature range (CSTR)s: (a) 18Mn-cast, (b) Ni-cast alloys.
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Fig. 4. The microstructure and fractured surface of the 18Mn-cast specimens on-heating hot ductility tested at: (a) and (d) 900 1C, (b) and (e) 1200 1C, (c) and (f) 1300 1C.
Fig. 5. The microstructure and fractured surface of the Ni-cast specimens on-heating hot ductility tested at: (a) and (d) 900 1C, (b) and (e) 950 1C, (c) and (f) 1200 1C, respectively.
ductility curves (Fig. 3), the ductility increased in the 18Mn alloy and the change of the ductility curvature was observed in the Niadded alloy after initiation of the dynamic recrystallization at 900 1C. Even though the onset temperature of dynamic recrystallization was 900 1C for both alloys, the degree of recrystallization was different. As shown in Figs. 4(a) and 5(a), there were more partially recrystallized grains in the 18Mn than the Ni-added alloy. Dynamic recrystallization occurs when the applied strain is greater than the critical strain. The Zener–Hollomon parameter (Z) which
describes the high temperature strain of a material: Z ¼ έ exp Q =RT
ð1Þ
where έ is the strain rate, Q is the activation energy for deformation, R is the gas constant, and T is the absolute temperature. As shown in Eq. (1), the critical strain for dynamic recrystallization depends on the deformation temperature, the strain rate and chemical composition dependent activation energy [31]. Therefore, the effect of grain size difference [30] and fine titanium
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Fig. 6. (a) M3C/γ eutectic at vicinity of the crack in Fig. 4(b) (open circle), (b) M3P/γ eutectic at vicinity of the crack in Fig. 5(b) (open circle), and (c) M3C/γ eutectic at vicinity of the crack in Fig. 5(c) (open circle).
Fig. 7. The microstructure and fractured surface of the on-cooling hot ductility tested 18Mn-cast specimens: (a) and (d) microstructure and fractured surface of 1200 1C, (b) microstructure of 1000 1C, (c) and (e) microstructure and fractured surface of 900 1C.
carbide particles (o0.5 μm) [32] of the Ni-added alloy can influence the onset temperature of the dynamic recrystallization and retard the growth of the nuclei of the dynamic recrystallization, are limited due to the slight difference in grain size and small quantities of TiC precipitates. Consequently, for the condition of the same test temperature and strain rate, the chemical composition of each material is the primary factor for the dynamic recrystallization. In view of the chemical composition difference between 18Mn and Ni-added alloys, the manganese and nickel contents varied as 18.5 wt% Mn and 0 wt% Ni in the 18Mn-cast specimen and less than 18 wt%Mn, and more than 5 wt%Ni in the Ni-cast specimen while the carbon content is same as 0.6 wt%. Cabaňas et al. [31] and Hamada et al. [33] reported that manganese delays dynamic recrystallization in the Fe–Mn binary alloys and high manganese TWIP steel due to the increase in the activation energy for hot working with increasing manganese content. In addition, solutes that increase the stacking fault energy (SFE) will favor dynamic recovery, while solutes that reduce the stacking fault energy will favor dynamic recrystallization as the primary softening mechanism. As a result of the thermo-dynamic calculation considering alloying elements and temperature using thermodynamic databases and equations [34–36], the stacking fault
energy of the Ni-added was greater than the 18Mn alloy with a difference of 4–5 mJ/m2 throughout the wide temperature range. Though no exact data was available for the high temperature stacking fault energy, it was expected to increase with temperature and the SFE difference between the 18Mn and the Ni-added alloy would be maintained. Consequently, the degree of dynamic recrystallization in the Ni-added alloy was smaller than the nickelfree 18Mn alloy due to the increase in stacking fault energy resulting in retarded dynamic recrystallization by the nickel addition. The flow stress level of two cast materials can also be explained by Mn–C dipoles short range ordering due to the high carbon content, as studied by Saeed-Akbari et al. [37]. The theoretical ordering index (TOI), which was defined for providing a measurement of the fraction of carbon and manganese atoms in different Fe–Mn–C systems [37], was 0.15 and 0.17 for the 18Mn and the Ni-added alloys, respectively. For TOI values between 0.1 and 0.3, the population of CMn6 octahedral clusters that can act as obstacles for dislocations increased. Therefore, as shown in Figs. 4(a), 7(b) and Figs. 5(a), 8(b), the dynamic recrystallization of both specimens were limited and only partial dynamic recrystallized grains were observed due to the high flow stress level of the two cast materials at a fast strain rate condition.
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Fig. 8. The microstructure and fractured surface of the on-cooling hot ductility tested Ni-cast specimens: (a) microstructure of 1200 1C, (b) microstructure of 1050 1C, (c) and (d) microstructure and fractured surface of 900 1C, respectively.
3.2.2. The liquation cracking As the temperature increased below the eutectic temperature of γ/M3C, the hot ductility of the 18Mn increased slightly, but the Ni-added alloy gradually decreased in the temperature range of 900–1000 1C even though dynamic recrystallization occurred after 900 1C. As shown in Figs. 5(b), 6(b), and 5(e), the melting trace was observed along grain boundaries and the fractured surface due to the (Fe, Mn)3P/γ eutectic melting occurred at 950 1C. Fig. 6(b) is the M3P/γ eutectic, with an enlarged image of the open circle shown in Fig. 5(b). Mintz et al. reported that the M3P/γ eutectic temperature can be lowered to 950 1C with the highly concentrated carbon at the interdendritic area and grain boundaries. Consequently, the ductility drop at 950 1C of the Ni-added alloy was due to the melting of the M3P/γ eutectic constituent at the grain boundaries that were formed in the as-cast condition. Consequently, in view of the hot ductility behavior, the nickel addition to the high manganese austenitic steel seems to have a detrimental effect of retarding dynamic recrystallization and the low melting eutectic constituent, M3P/γ formation by reducing the solubility of phosphorus in the austenite matrix. As the temperature increases to the M3C/γ eutectic temperature of the each specimen (Fig. 3), the ductility abruptly decreases by eutectic melting of M3C/γ, resulting in the formation of the liquid film on the grain boundaries that accelerates the change in fracture mode from transgranular to intergranular as shown in Figs. 4(b), 6(a), 4 (e), and Figs. 5(c), 6(c), 5(f). Fig. 6(a) and (c) are M3C/γ eutectics, enlarged images of open circles in Figs. 4(b) and 5(c), respectively. As the temperature approaches the zero ductility temperature (ZDT), the degree of ductility drop of both specimens gently decreases while the liquid penetrated and wetted all over the grain boundaries as shown in Fig. 4(c) and (f) at 1300 1C (ZDT). At this stage, no dynamic recrystallization occurred because the
fracture strain was lower than the critical strain for the initiation of recrystallization. The ductility of this stage is a function of the extent of liquid formation via incipient melting at Si, P and S segregated grain boundaries and liquid continuity along the grain boundaries [29].
3.2.3. The on-cooling behavior On-cooling ductility recovery of both alloys started under 1100 1C near the eutectic temperature of γ/M3C. Figs. 7(a), (d), and 8(a) are typical microstructures near the fractured region and fracture surface of 18Mn and Ni-added alloys tested on-cooling at 1200 1C, respectively. The microstructures of both specimens were similar to those in the ductility gently decreased stage of the onheating test. The difference is the wider grain boundaries due to the prolonged exposure near the ZDT, resulting in extensive liquid formation along the grain boundaries. Even though most of the melted grain boundaries have started to solidify at this temperature range under the equilibrium solidus temperature of both alloys (Fig. 2), but as shown in the fractography at 1200 1C in Fig. 7(d), boundaries were penetrated by the M3C/γ eutectic liquid in a wetted condition at the moment of fracture. With a further decrease in temperature below the eutectic temperature of M3C/γ, the ductility increased due to the depletion of the liquid along the grain boundaries with the prevention of the initial crack by recrystallized grains and grain boundary migration (Figs. 3, 8(a), and (b), microstructures of the on-cooling tested 18Mn at 1000 1C and Ni-added alloy at 1050 1C, respectively.). As shown in Fig. 3, a ductility drop was observed in the temperature range of 1000 1C to 900 1C for the 18Mn and 1050 1C to 900 1C for the Ni-added alloy, respectively. When compared to microstructures of the on-heating tested specimens, preciously melted and resolidified “ghost
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boundaries [29]” were clearly observed as shown in Figs. 7(c) and 8 (c). Authors also mentioned that the previously melted and resolidified boundaries were enriched in solute elements, including P, S and Si, compared to newly formed boundaries resulting in significant embrittlement, which are favored for crack initiation and propagation along the boundaries. No recrystallized grains and relatively long and straight cracks along the previously melted and resolidified grains were observed in the on-cooling both alloys at 900 1C (Figs. 7(c), 8(c)). The fractographies also show no liquidrelated evidence, but the completely intergranular solid-state fracture along resolidified boundaries with some proof, i.e. a striation or ripped eutectics on the fracture (Figs. 7(e) and 8(d)). This ductility dip phenomena can be influenced by the peak temperature of the on-cooling test in terms of the eutectic melting, quantities of the liquid formation, degree of grain boundary liquid penetration, the grain boundary migration and the final grain size that can affect the dynamic recrystallization. With the on-cooling temperature of ZDT, no dynamic recrystallization or ductility dip phenomena of either cast alloy was observed near at 900 1C as seen in the multi-pass weld heat affected zone of the Ni added aswelded specimen in this study. If the peak temperature increases to the zero strength temperature (ZST), the recrystallization temperature can be shifted to a higher temperature, and vice versa. This is primarily due to the effect of the grain size on dynamic recrystallization and the quantities of eutectic constituents and possible precipitation that can affect the embrittlement or hardening of the matrix.
3.3. Results of the Varestraint test Fig. 9 shows the results of the Varestraint tests for the 18Mn and the Ni-added alloys. As shown in Figs. 10 and 11, all of cracks were counted and measured at the HAZ of the autogeneous gas tungsten arc welded 18Mn and Ni-added alloy. The total crack length (TCL) and maximum crack length (MCL) those were measured in direction perpendicular to the fusion line of Niadded alloy were slightly larger than those of the 18Mn, and the TCL and MCL of 18Mn alloys were nearly the same as the TCL and MCL of 18Mn–0.6C weld metal HAZ with the 4% augmented strain of the previous study [23]. Besides the conventional TCL and MCL measurements, the other criterion designed by Lee [38], termed “cracked HAZ length (CHL)”, measured along the fusion line in the HAZ where cracks formed was also measured to directly compare the hot ductility test evaluating HAZ cracking susceptibility along the same peak temperature region parallel to the fusion line. The CHL, the length parallel to the fusion line between the start point of the cooling portion of the weld pool and the region where the
Fig. 10. Cracked HAZ length (CHL) of 18Mn-cast and Ni-cast Varestraint tested specimens.
last crack was observed, of the 18Mn and the Ni-added alloy were 3.6 and 5.5 mm, respectively (Figs. 10, 11). Based on the directly-measured, on-cooling hot crack susceptibility test for the actual welds, the CHL results can be another good indicator for the evaluation of specimens with similar conventional criterions. Fig. 8 shows microstructures of the Varestraint tested specimens. Isothermal lines from the fusion line are indicated in Fig. 11(a) and (d) as the calculation results of the three dimensional Rosenthal's equation. As shown in these figures, cracks in the HAZ were initiated from the fusion line and propagated towards the base metal until the lowest eutectic temperature of each specimen, 1080 1C of M3C/γ and 950 1C of M3P/γ for the 18Mn and Ni-added alloys, respectively. Fig. 11(b) and (e) are the enlarged images of the M3C/γ eutectic for the 18Mn and Ni-added alloys, while Fig. 11(f) is the M3P/γ eutectics near the end tip of the crack in the Ni-added alloy. The ductility dip crack along the migrated grain boundaries were also observed below the lowest eutectic temperature as shown in Fig. 11(c) and (d). Consequently, cracks initiated above the solidus or eutectic temperature by a liquation mechanism and propagated along the liquid-penetrated grain boundaries followed the solid-state crack formation due to the resolidified and embrittled boundaries with the same crack formation mechanism of the as-welded alloys.
4. Conclusions The hot cracking susceptibility of two as-cast alloys with different manganese and nickel contents were investigated by the hot ductility and the Varestraint tests for clarifying the mechanism of cracks in a multi-pass weld heat affected zone of the cryogenic high manganese steel. Conclusions of this study are as follows.
Fig. 9. Results of the Varestraint test: total crack length, maximum crack length, cracked HAZ length.
1. The additional nickel on fully austenitic high alloyed manganese steels can significantly reduce the phosphorus solubility in the matrix compared to nickel-free austenitic high manganese steel, resulting in more severe phosphorus segregation in the grain or sub grain boundaries and M3P/γ eutectic formation during solidification in the as-cast condition.
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Fig. 11. Microstructure of the Varestraint tested specimens: (a) low magnification image of the fusion zone and heat affected zone of 18Mn-cast specimen, (b) enlarged image of the open square area in Fig. (a), (c) crack across the line of Te along migrated grain boundary of 18Mn-cast specimen weld HAZ, (d) low magnification image of the fusion zone and heat affected zone of the Ni-cast specimen, (e) and (f) enlarged image of the solid and open square area in Fig. (d), respectively.
2. Hot cracking susceptibility criteria from the hot ductility test, BTR and CSTR of 18Mn and Ni-added cast alloys were similar but the overall ductility of the Ni-added alloy was less than that of the 18Mn alloy due to the low temperature of eutectic melting of the γ/M3P, coarse and randomly distributed MnS inclusions, and the low degree of the dynamic recrystallization by the higher stacking fault energy compared to the 18Mn alloy. 3. Dynamic recrystallization was limited due to the high carbon contents of both 18Mn and Ni-added alloys resulting in only partial dynamic recrystallization under a condition of fast strain rate. 4. Significant ductility drops were observed in both 18Mn and Ni-added alloys as a result of M3C/γ eutectic melting at each eutectic temperatures during on-heating tests. That was attributed to a change of fracture mode from transgranular to intergranular and was accelerated by the liquid films formation along grain boundaries. 5. With the on-cooling peak temperature of ZDT, no dynamic recrystallization following the ductility dip phenomena of both cast alloys were observed near 900 1C as observed in the multipass weld heat affected zone of the Ni added welds. 6. Total and maximum crack lengths and the cracked HAZ length that allow direct measurements of the on-cooling hot crack susceptibility of the actual welds for the Ni-added alloy was bigger than the 18Mn alloy as a result of the Varestraint test.
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