metals Article
Influence of Post-Weld Heat Treatment on the Microstructure, Microhardness, and Toughness of a Weld Metal for Hot Bend Xiu-Lin Han 1 , Da-Yong Wu 2 , Xiang-Ling Min 1 , Xu Wang 1 , Bo Liao 2 and Fu-Ren Xiao 2, * 1 2
*
CNPC Bohai Petroleum Equipment Manufacture Co., Ltd., Qingxian 062658, China;
[email protected] (X.-L.H.);
[email protected] (X.-L.M.);
[email protected] (X.W.) Key Laboratory of Metastable Materials Science & Technology, College of Materials Science & Engineering, Yanshan University, Qinhuangdao 066004, China;
[email protected] (D.-Y.W.);
[email protected] (B.L.) Correspondence:
[email protected]; Tel./Fax: +86-335-807-7110
Academic Editor: Hugo F. Lopez Received: 31 December 2015; Accepted: 18 March 2016; Published: 25 March 2016
Abstract: In this work, a weld metal in K65 pipeline steel pipe has been processed through self-designed post-weld heat treatments including reheating and tempering associated with hot bending. The microstructures and the corresponding toughness and microhardness of the weld metal subjected to the post-weld heat treatments have been investigated. Results show that with the increase in reheating temperature, austenite grain size increases and the main microstructures transition from fine polygonal ferrite (PF) to granular bainitic ferrite (GB). The density of the high angle boundary decreases at higher reheating temperature, leading to a loss of impact toughness. Lots of martensite/austenite (M/A) constituents are observed after reheating, and to a large extent transform into cementite after further tempering. At high reheating temperatures, the increased hardenability promotes the formation of large quantities of M/A constituents. After tempering, the cementite particles become denser and coarser, which considerably deteriorates the impact toughness. Additionally, microhardness has a good linear relation with the mean equivalent diameter of ferrite grain with a low boundary tolerance angle (2˝ ´8˝ ), which shows that the hardness is controlled by low misorientation grain boundaries for the weld metal. Keywords: M/A constituent; cementite; toughness; microhardness; EBSD
1. Introduction Pipeline steels used to transport crude oil or natural gas over long distances under high pressure essentially require extraordinary strength and toughness; thus, numerous studies have been conducted to achieve this purpose. The acicular ferrite (AF) microstructure in pipeline steel, which has been considered to have the optimum mechanical properties, has become the hotspot of study. Xiao et al. [1] reported that the ferrite microstructures in pipeline steel were divided into polygonal ferrite (PF), quasi-polygonal ferrite (QF) or massive ferrite (MF), granular bainite (GB), and bainitic ferrite (BF), according to the metallographical identification and classification of the phases. Based on this work, Wang et al. [2,3] claimed that AF was a complex consisting of QF, GB, and BF, with dispersed islands of second phases in the matrix. The higher dislocation density and sub-boundaries within QF and GB provided greater strength for the pipeline steel, compared to the conventional PF-pearlite structure. Similar work by Chung et al. [4] showed that the strengthening effect can be enhanced by fine ferrite laths, precipitates, and martensite/austenite (M/A) constituents with small island and thin film inside the ferrite structure.
Metals 2016, 6, 75; doi:10.3390/met6040075
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Currently, thermo-mechanical processing is considered as the primary route for developing American Petroleum Institute (API) grade pipeline steels because such a process definitively influences the final microstructure. Most studies have focused on the effect of processing parameters on the Metals 2016, 6, 75 2 of 11 microstructure of the pipe body and/or the welding heat-affected zone. However, steel pipe products include not only the straight pipes but also a number of accessories, such as bends. Post-weld heat Currently, thermo‐mechanical processing is considered as the primary route for developing treatments including reheating and(API) tempering with hot bending for the welds in pipeline American Petroleum Institute grade associated pipeline steels because such a process definitively steels have not received sufficient attention. In particular, limited information is available influences the final microstructure. Most studies have focused on the effect of processing about parameters on the microstructure of the pipe body and/or the welding heat‐affected zone. the effect of the post-weld heat treatments associated with hot bending on the microstructures and However, steel pipe products include not only the straight pipes but also a number of accessories, properties of the weld joint. The bends have become a key part of the pipeline project. Therefore, it is such as bends. Post‐weld heat treatments including reheating and tempering associated with hot necessary to explore the evolution of the microstructure and related properties of the weld metal in bending for the welds in pipeline steels have not received sufficient attention. In particular, limited pipes during hot bending. information is available about the effect of the post‐weld heat treatments associated with hot Additionally, to explore the relationship between microstructure and properties, numerous bending on the microstructures and properties of the weld joint. The bends have become a key part researchers consider the orientation of the ferrite grains with respect to each other via EBSD technology. of the pipeline project. Therefore, it is necessary to explore the evolution of the microstructure and related properties of the weld metal in pipes during hot bending. For example, in the analysis of toughness, Díaz-Fuentes et al. [5] reported that the microstructural unit Additionally, to explore the relationship between microstructure and defined properties, controlling brittle fracture in an acicular ferrite microstructure could be as anumerous set of adjacent researchers consider the orientation of the ferrite grains with respect to ˝ each other via EBSD ferrite plates with crystallographic misorientations below an angle of 15 . For the strength analysis, technology. For example, in the analysis of toughness, Díaz‐Fuentes et al. [5] reported that the Iza-Mendia and Gutiérrez [6] revealed that the optical ferrite mean linear intercept was equivalent to microstructural unit controlling brittle fracture in an acicular ferrite microstructure could be ˝ . In this the mean equivalent diameter (MED) computed EBSD scansmisorientations using a threshold of 2of defined as a set of adjacent ferrite plates with from crystallographic below angle an angle work, the as-received material is a Iza‐Mendia and weld metal in Gutiérrez K65 pipeline steel. After post-weld heatmean treatments 15°. For the strength analysis, [6] revealed that the optical ferrite linear intercept was equivalent to the mean equivalent diameter (MED) computed from EBSD scans associated with hot bending, the relationship between microstructure and hardness was analyzed using a threshold angle of 2°. In this work, the as‐received material is a weld metal in K65 pipeline via EBSD. steel. After post‐weld heat treatments associated with hot bending, the relationship between During the actual hot bending process, the weld metal is usually set up at the top of the pipe, so microstructure and hardness was analyzed via EBSD. the plastic deformation can be neglected [7,8]. Therefore, the hot bending process can be simulated by During the actual hot bending process, the weld metal is usually set up at the top of the pipe, a reheating treatment. Our previous showed thatTherefore, increasing ofhot Ni bending content in weldcan metals so the plastic deformation can work be neglected [7,8]. the process be could by a reheating treatment. Our previous work showed that increasing of Ni content in improvesimulated the mechanical properties of the weld metals after simulated post-weld heat treatments for weld metals could improve the mechanical properties of the weld metals after simulated post‐weld hot bends [9], while the effects of post-weld heat treatments on the microstructure and mechanical heat treatments for hot bends [9], while the effects of post‐weld heat treatments on the properties have not sufficiently investigated. In this paper, the effects of post-weld reheating and microstructure and mechanical properties have not sufficiently investigated. In this paper, the tempering temperature on the microstructure, toughness, and hardness of the weld metal with high effects of post‐weld reheating and tempering temperature on the microstructure, toughness, and Ni content were investigated. hardness of the weld metal with high Ni content were investigated. 2. Experimental Materials and Procedure 2. Experimental Materials and Procedure 2.1. Materials and the Post‐Weld Heat Treatments 2.1. Materials and the Post-Weld Heat Treatments The base metal was a commercial K65 pipeline steel with 1219 mm outside diameter and 30.8 The base metal was a commercial K65 pipeline steel with 1219 mm outside diameter and mm wall thickness. The double face four wire tandem submerged arc welding was performed 30.8 mm wall thickness. The double face four wire tandem submerged arc welding was performed according to the actual welding process for a K65 hot bend. Test pieces of weld metal with the size according to the actual welding process for a K65 hot bend. Test pieces of weld metal with the size 10 mm × 10 mm × 80 mm were cut from the outer side of the weld metal transversely to the weld 10 mm ˆ 10 mm ˆ 80in mm were the outer side of of thethe weld metal transversely to the weld line, line, as shown Figure 1. cut The from chemical compositions out‐side weld metal are shown in as shown in Figure 1. The chemical compositions of the out-side weld metal are shown in Table 1. Table 1.
Figure 1. Schematic diagram for sampling method of the impact specimens (WZ: weld zone; HAZ: heat affected zone; BM: base metal).
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Figure 1. Schematic diagram for sampling method of the impact specimens (WZ: weld zone; HAZ: Metals heat affected zone; BM: base metal). 2016, 6, 75
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Table 1. Chemical compositions of the weld metal (wt. %). Table 1. Chemical compositions of the weld metal (wt. %).
Elements C Mn Si P S Mo Ni Cr Cu V Nb Ti Al N B Ceq 1 Pcm 2 Mn Si P S Mo Ni Cr Cu V Ti N B Elements C Nb Al Ceq 1 Pcm 2 Amount 0.06 1.67 0.28 0.019 0.005 0.161 0.52 0.123 0.105 0.01 0.036 0.015 0.019 0.008 0.0008 0.44 0.18 Amount
0.06
1.67
1 1Ceq Ceq ` C “C
0.28
0.019
0.005
0.161
0.52
0.123
0.105
0.01
0.036
0.015
0.019
0.008
0.0008
0.44
Si Mn Cu Cr Ni Mo Mo V Mn Ni Cu Cr Mo V 2 Pcm ` C Si ` Mn ` Cu ` Cr ` Ni ` V `5B Mn ` Ni ` Cu `Cr ` Mo ` V ; 2; Pcm “C ` 5B 66 1515 55 3030 2020 60 10 60 15 10
0.18
Post‐weld heat heat treatments treatments were conducted a Gleeble‐3500 thermal simulator New Post-weld were conducted on on a Gleeble-3500 thermal simulator (DSI, (DSI, New York, York, NY, USA), and the processing schedules are provided in Figure 2. The time from room NY, USA), and the processing schedules are provided in Figure 2. The time from room temperature to temperature to reheating temperature was 60 s, and the temperatures ranged from 990 °C to 1080 reheating temperature was 60 s, and the temperatures ranged from 990 ˝ C to 1080 ˝ C. After maintaining °C. After maintaining the reheating temperature for 60 s, some quenched samples were obtained for the reheating temperature for 60 s, some quenched samples were obtained for prior austenite grain prior austenite grain size observation and others were cooled to room temperature at a cooling rate size observation and others were cooled to room temperature at a cooling rate of 10 ˝ C/s. Finally of 10 °C/s. Finally the samples were tempered 550 °C and to 680 °C, respectively, to simulate the ˝ C and the samples were tempered at 550 680 ˝ C,at respectively, simulate the post-weld treatments post‐weld treatments associated with an industrial hot bending production process as closely as associated with an industrial hot bending production process as closely as possible. possible.
Figure 2. Schematic illustration of post‐weld heat treatment schedule for experimental weld metal. Figure 2. Schematic illustration of post-weld heat treatment schedule for experimental weld metal.
2.2. Mechanical Properties 2.2. Mechanical Properties The low-temperature low‐temperature impact impact toughness toughness and microhardness of of the the specimens specimens subjected to the the The and microhardness subjected to reheating and tempering processes were measured. Charpy impact tests were performed at −40 °C reheating and tempering processes were measured. Charpy impact tests were performed at ´40 ˝ C using the Charpy V‐notch (CVN) specimens in accordance with the standard of ASTME23‐02. The using the Charpy V-notch (CVN) specimens in accordance with the standard of ASTME23-02. The V‐notch was was machined at the center of weld the metal weld in metal in the cross ofsection of the The V-notch machined at the center of the the cross section the weld. The weld. toughness toughness values were calculated as an average of three measurements. The Vickers microhardness values were calculated as an average of three measurements. The Vickers microhardness was measured was measured a FM‐ARS 9000 hardness testing machine Tokyo, using a FM-ARSusing 9000 hardness testing machine (Future-Tech, Tokyo,(Future‐Tech, Japan) at a load of 500Japan) gf, andat thea load of 500 gf, and the indentation area was selected randomly in the weld metal. The hardness test indentation area was selected randomly in the weld metal. The hardness test results represented the results represented the average of five measurements. average of five measurements. 2.3. Microstructural Characterization 2.3. Microstructural Characterization The CVN CVN specimen specimen fracture fracture surface surface was was observed observed by by using using aa KYKY-2800 KYKY‐2800 scanning scanning electron electron The microscope (SEM) (SEM) (KYKY, Beijing, China) to examine the cleavage fracture and the crack microscope (KYKY, Beijing, China) to examine the cleavage fracture and the crack propagation propagation path. Specimens from the transverse cross section of the steels were path. Specimens derived fromderived the transverse cross section planes of the planes steels were mechanically mechanically polished and etched by a 3% nital solution. The microstructures were then observed polished and etched by a 3% nital solution. The microstructures were then observed by using a by using a Hitachi S‐4800 SEM (Hitachi, Tokyo, Japan). The initial austenite grain boundaries were Hitachi S-4800 SEM (Hitachi, Tokyo, Japan). The initial austenite grain boundaries were observed on observed on quenched specimens etched in an aqueous picric Toacid solution. To observe the M/A quenched specimens etched in an aqueous picric acid solution. observe the M/A constituent, the constituent, the LePera reagent of 1 g sodium metabisulfite in 100 mL of distilled water, mixed with LePera reagent of 1 g sodium metabisulfite in 100 mL of distilled water, mixed with 4 g of picric acid dissolved in 100 mL of alcohol was used. The M/A constituents were identified by optical microscopy (OM, Zeiss, Oberkochen, Germany). Samples for the electron backscatter diffraction (EBSD) analyzer
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(EDAX, Mahwah, NJ, USA) were electrolytically polished in a solution of 8% perchloric acid and 92% ethanol. The EBSD maps were analyzed by Channel 5 software (TSL OIM 7, EDAX, Mahwah, NJ, USA, 2011) from Oxford-HKL. Thin foils for transmission electron microscopy (TEM) were prepared by mechanical thinning from 300 µm to 50 µm, along with twin-jet electropolishing in an electrolyte of 8% perchloric acid and 92% ethanol; and the examinations were performed using a JEM-2010 TEM (JEOL, Tokyo, Japan) with 200 kV. 3. Results 3.1. Properties Charpy impact energy (´40 ˝ C) and Vicker’s hardness results of the samples under different conditions are listed in Table 2. The tempered specimens reheated at 990 ˝ C and 1030 ˝ C achieved satisfactory impact toughness energy within 86 J to 133 J. The values of impact energy, 24 J to 29 J, were significantly lower for the tempered specimens reheating at 1080 ˝ C. The microhardness of the tempered specimens reheated at 990 ˝ C and 1030 ˝ C was higher than that of the tempered specimens reheated at 1080 ˝ C. Table 2. Toughness and microhardness of the weld metal. Reheating Temperature (˝ C)
990
1030
1080
Tempering Temperature (˝ C)
Charpy Impact Energy (´40 ˝ C, J) Discrete Values
Average Value
Microhardness (HV)
Error Range
Discrete Values
Average Value
Error Range
550
115, 152, 126
131
˘21
248, 257, 257, 246, 253
252
˘6
680
109, 132, 146
129
˘20
249, 254, 252, 230, 250
247
˘17
550
138, 143, 118
133
˘15
268, 250, 256, 259, 249
256
˘12
680
96, 92, 70
86
˘16
255, 259, 259, 254, 243
254
˘11
550
45, 18, 24
29
˘16
230, 238, 244, 229, 232
235
˘9
680
16, 19, 37
24
˘13
229, 238, 235, 235, 220
231
˘11
3.2. Microstructure The prior austenite grain and microstructures after reheating at different temperatures are shown in Figure 3. The grain size of austenite varied insignificantly at 990 ˝ C and 1030 ˝ C, while it increased markedly at 1080 ˝ C (Figure 3a–c). This was associated with the quick dissolution of fine Nb precipitates at a higher reheating temperature [10]. The prior austenite microstructures were also affected on the cooled microstructures. When the reheating temperature was 990 ˝ C, the precipitates hardly dissolved, the austenite grain size was relatively small, and the alloy elements distribution was heterogeneous, so the phase transformation temperature was higher. The main microstructures were PF and QF (Figure 3d). A high austenitizing temperature could improve austenite’s homogeneity and hardenability; otherwise it is prone to form low-temperature microstructures. Therefore, a mixed microstructure of fine QF and GB was obtained at 1030 ˝ C (Figure 3e), and the microstructure changed to mainly coarse GB at 1080 ˝ C (Figure 3f). Additionally, as the ferrite nucleated and grew, partitioning of carbon took place at the austenite-ferrite interface; the austenite-enriched carbon would transform to its M/A constituent. After tempering, the PF and QF showed little change, while the GB and M/A constituent changed remarkably because the recovery and recrystallization occurred in parts of GB, resulting in smaller PF grains distributed at the lath ferrite boundary [11,12], as clearly illustrated in Figure 4. Wang et al. [3]
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claimed that PF generally grows into equiaxed grains with smooth and continuous boundaries. However, QF boundaries are irregular and jagged, even discontinuous. BF consists of many elongated parallel ferritic laths. Compared to BF, GB contains irregular ferritic laths with no clear boundaries. According to different microstructural characteristics, different ferrites of PF, QF, and GB have been marked. In addition, degenerate pearlite was observed after tempering at 550 ˝ C, as is shown in Figure 5a. There were two kinds of particles observed in the SEM micrographs: one was proven to be Ti-oxide in the weld metal (Figure 5b), randomly distributed in the microstructure, which could serve as the nucleation site of intra-granular ferrite; another kind, mainly distributed on the grain boundaries, was identified as carbides, produced by the tempering transformation of the M/A constituents (Figure 5c). With increasing tempering temperatures, the degenerate pearlite disappeared while the amount of carbides allocated on the grain boundaries increased. The evolution of the M/A constituent through reheating at different temperatures, before and ˝ C, is presented in Figure 6. Before tempering, a large amount of M/A islands afterFigure tempering at 550 3. Optical images of prior austenite grain boundary (a–c) and microstructures (d–f) at weredifferent reheating temperatures: (a,d) 990 °C; (b,e) 1030 °C; (c,f) 1080 °C. dispersed in the ferrite matrix when reheating at 990 ˝ C and 1030 ˝ C. However, for the specimen reheating at 1080 ˝ C, the M/A constituent not only distributed in the ferritic matrix, but also segregated on the priortempering, austenite grain tempering, the amount of M/A islands sharply After the boundaries. PF and QF After showed little change, while the GB and decreased M/A constituent and the retained M/A islands were refined. changed remarkably because the recovery and recrystallization occurred in parts of GB, resulting in Metals 2016, 6, 75 5 of 11 smaller PF grains distributed at the lath ferrite boundary [11,12], as clearly illustrated in Figure 4. Wang et al. [3] claimed that PF generally grows into equiaxed grains with smooth and continuous boundaries. However, QF boundaries are irregular and jagged, even discontinuous. BF consists of many elongated parallel ferritic laths. Compared to BF, GB contains irregular ferritic laths with no clear boundaries. According to different microstructural characteristics, different ferrites of PF, QF, and GB have been marked. In addition, degenerate pearlite was observed after tempering at 550 °C, as is shown in Figure 5a. There were two kinds of particles observed in the SEM micrographs: one was proven to be Ti‐oxide in the weld metal (Figure 5b), randomly distributed in the microstructure, which could serve as the nucleation site of intra‐granular ferrite; another kind, mainly distributed on the grain boundaries, was identified as carbides, produced by the tempering transformation of the M/A constituents (Figure 5c). With increasing tempering temperatures, the degenerate pearlite disappeared while the amount of carbides allocated on the grain boundaries increased. The evolution of the M/A constituent through reheating at different temperatures, before and after tempering at 550 °C, is presented in Figure 6. Before tempering, a large amount of M/A islands were dispersed in the ferrite matrix when reheating at 990 °C and 1030 °C. However, for the Figure 3.3. Optical Optical images of prior austenite boundary (a–c) and microstructures (d–f) at Figure images of prior austenite graingrain boundary (a–c) and microstructures (d–f) at different specimen reheating at 1080 °C, the M/A constituent not only distributed in the ferritic matrix, but ˝ C; (b,e) 1030 ˝ C; (c,f) 1080 ˝ C. different reheating temperatures: (a,d) 990 °C; (b,e) 1030 °C; (c,f) 1080 °C. reheating temperatures: (a,d) 990 also segregated on the prior austenite grain boundaries. After tempering, the amount of M/A islands decreased sharply and the retained M/A islands were refined. After tempering, the PF and QF showed little change, while the GB and M/A constituent changed remarkably because the recovery and recrystallization occurred in parts of GB, resulting in smaller PF grains distributed at the lath ferrite boundary [11,12], as clearly illustrated in Figure 4. Wang et al. [3] claimed that PF generally grows into equiaxed grains with smooth and continuous boundaries. However, QF boundaries are irregular and jagged, even discontinuous. BF consists of many elongated parallel ferritic laths. Compared to BF, GB contains irregular ferritic laths with no clear boundaries. According to different microstructural characteristics, different ferrites of PF, QF, and GB have been marked. In addition, degenerate pearlite was observed after tempering at 550 °C, as is shown in Figure 5a. There were two kinds of particles observed in the SEM micrographs: one was proven to be Ti‐oxide in the weld metal (Figure 5b), randomly distributed in the microstructure, which could serve as the nucleation site of intra‐granular ferrite; another kind, mainly distributed on the grain boundaries, was identified as carbides, produced by the tempering transformation of the M/A constituents (Figure 5c). With increasing tempering temperatures, the 4. Cont. degenerate pearlite disappeared while the Figure amount of carbides allocated on the grain boundaries increased. The evolution of the M/A constituent through reheating at different temperatures, before and after tempering at 550 °C, is presented in Figure 6. Before tempering, a large amount of M/A islands were dispersed in the ferrite matrix when reheating at 990 °C and 1030 °C. However, for the specimen reheating at 1080 °C, the M/A constituent not only distributed in the ferritic matrix, but also segregated on the prior austenite grain boundaries. After tempering, the amount of M/A
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Figure 4. SEM micrograph for specimens treated under different conditions: (a) Prior to treatments; Figure 4. SEM micrograph for specimens treated under different conditions: (a) Prior to treatments; Figure 4. SEM micrograph for specimens treated under different conditions: (a) Prior to treatments; (b) RT = 990 °C, TT = 550 °C; (c) RT = 990 °C, TT = 680 °C; (d) RT = 1030 °C, TT = 550 °C; (e) RT = 1030 (b) RT = 990 ˝ C, TT = 550 ˝ C; (c) RT = 990 ˝ C, TT = 680 ˝ C; (d) RT = 1030 ˝ C, TT = 550 ˝ C; (b) RT = 990 °C, TT = 550 °C; (c) RT = 990 °C, TT = 680 °C; (d) RT = 1030 °C, TT = 550 °C; (e) RT = 1030 ˝ C; ˝ C; ˝ C; °C, ˝ C, TT TT = ˝680 °C; =(f) RT = 1080 °C; = TT 1080 = 550 °C; TT (g) = RT 550 = 1080 °C. (RT: reheating (e) RT°C, = 1030 C, TT 680 (f) RT (g) TT RT= =680 1080 = 680 ˝ C. TT = 680 °C; (f) RT = 1080 °C; TT = 550 °C; (g) RT = 1080 °C, TT = 680 °C. (RT: reheating °C, temperature; TT: tempering temperature). (RT: reheating temperature; TT: tempering temperature). temperature; TT: tempering temperature).
Figure 5. Cont.
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Figure 5. SEM micrographs for degenerate pearlite (a); intra‐granular ferrite nucleated around Figure 5. SEM micrographs for degenerate pearlite (a); intra‐granular ferrite nucleated around Figure 5. SEM micrographs for degenerate pearlite (a); intra-granular ferrite nucleated around Ti-oxide Ti‐oxide particle (insert: EDX result of Ti‐oxide particle) (b), and carbides (c) in specimens reheated Ti‐oxide particle (insert: EDX result of Ti‐oxide particle) (b), and carbides (c) in specimens reheated particle (insert: EDX result of Ti-oxide particle) (b), and carbides (c) in specimens reheated at 990 ˝ C; at 990 °C; tempering at 550 °C (a,b) and 680 °C (c). at 990 °C; tempering at 550 °C (a,b) and 680 °C (c). tempering at 550 ˝ C (a,b) and 680 ˝ C (c).
Figure 6. M/A constituent observations of the samples reheated at different temperatures (990 °C (a); 1030 (b); 1080 (c)) before tempering and after tempering at 550 °C (reheated at 990 °C (d); 1030 ˝ C (a); Figure 6. samples reheated at at different temperatures (990(990 Figure 6. M/A M/A constituent constituent observations observations ofof the the samples reheated different temperatures °C °C (e); 1080 °C (f)). ˝ ˝ ˝
1030 (b); 1080 (c)) before tempering and after tempering at 550 C (reheated at 990 C (d); 1030 C (e); (a); 1030 (b); 1080 (c)) before tempering and after tempering at 550 °C (reheated at 990 °C (d); 1030 ˝ C (f)). 1080 3.3. Particles °C (e); 1080 °C (f)).
Multiple‐scale particles were observed via TEM. One of them with approximately 0.5 μm to 1.5
3.3. Particles 3.3. Particles μm was identified as Ti‐oxide (Figure 7a). A mass of intra‐granular ferrite nucleated on Ti‐oxides
was observed. particles Early work by observed Shim et al. examined the potency of various inclusions and Multiple-scale were via[13] TEM. One of them with approximately 0.5 µm to 1.5 µm Multiple‐scale particles were observed via TEM. One of them with approximately 0.5 μm to 1.5 precipitates, e.g., MnS, Ti 2O3, VN, TiMn2O3, NbC, VC, AlN for the nucleation of intra‐granular was identified as Ti-oxide (Figure 7a). A mass of intra-granular ferrite nucleated on Ti-oxides was μm was identified as Ti‐oxide (Figure 7a). A mass of intra‐granular ferrite nucleated on Ti‐oxides ferrite, and found that TiMn 2O3 particles showed great efficacy in inducing intra‐granular ferrite. observed. Early Early work by Shim al. [13]et examined the potency ofpotency various inclusions precipitates, was observed. work by etShim al. [13] examined the of various and inclusions and The addition of Ti2O3 nanoparticles to the weld metal was required in order to increase the density e.g., MnS, Ti O , VN, TiMn O , NbC, VC, AlN for the nucleation of intra-granular ferrite, and found precipitates, 2 e.g., 3 MnS, Ti22O3, 3 VN, TiMn2O3, NbC, VC, AlN for the nucleation of intra‐granular of the ferrite nucleation site. Therefore, a high density of fine ferrite was achieved, which was that TiMn Ofound particles showed efficacy in inducing ferrite. intra‐granular The addition offerrite. Ti2 O3 ferrite, and that TiMn 2Ogreat 3 particles showed great intra-granular efficacy in inducing 2 3 beneficial to the improvement of strength and toughness. The second particle, with a size of nanoparticles to the weld was in order to increase the density of the ferrite nucleation The addition of Ti 2O 3 nanoparticles to the weld metal was required in order to increase the density approximately 200 nm metal to 500 nm, required was a carbide identified as cementite (Figure 7b). The coarse site. Therefore, a high density of fine ferrite was achieved, which was beneficial to the improvement of the ferrite nucleation site. Therefore, a high density of fine ferrite was achieved, which was cementite particles embellishing the prior austenite grain boundaries resulted in the deterioration of of strength and toughness. The second particle, with a size of approximately 200 nm with to 500a nm, was toughness, because stress concentration occurred at the boundaries during the impact process, and beneficial to the improvement of strength and toughness. The second particle, size of the coarse cementite particles easily crack initiations an applied stress, which a carbide identified as cementite (Figure 7b). The coarse cementite particles embellishing thecoarse prior approximately 200 nm to 500 nm, was became a carbide identified as under cementite (Figure 7b). The eventually the impact toughness. The third particle type with a because size of approximately 50 austenite grainreduced boundaries resulted in the deterioration of toughness, stress concentration cementite particles embellishing the prior austenite grain boundaries resulted in the deterioration of nm to 200 nm was alloyed precipitates, including NbC, VC, AlN, and composite precipitates occurred at the boundaries during the impact process, and the coarse cementite particles easily became toughness, because stress concentration occurred at the boundaries during the impact process, and (Figure 7c). crack initiations under particles an applied stress,became which eventually reducedunder the impact toughness. Thewhich third the coarse cementite easily crack initiations an applied stress, particle typereduced with a size of approximately 50The nmthird particle to 200 nm wastype with alloyed precipitates, including NbC, eventually the impact toughness. a size of approximately 50 VC, and composite precipitates (Figureincluding 7c). nm AlN, to 200 nm was alloyed precipitates, NbC, VC, AlN, and composite precipitates (Figure 7c).
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Figure 7. TEM images thethree three scales scales particles particles inin specimens: (a)(a) Ti-oxide; (b) (b) cementite Figure 7. TEM images for for the specimens: Ti‐oxide; cementite decomposed from M/A island (insert: diffraction pattern); (c) MC carbides precipitated in matrix decomposed from M/A island (insert: diffraction pattern); (c) MC carbides precipitated in matrix (insert: diffraction pattern). (insert: diffraction pattern).
3.4. EBSD Analysis of the Microstructures
3.4. EBSD Analysis of the Microstructures
Figure 8 shows the inverse pole figures (IPF) of the ferrite phase with boundary misorientation
Figure 8 shows the inverse pole figures (IPF) of the ferrite phase with boundary misorientation maps obtained from EBSD measurement of the specimens after the whole heat treatments. ˝ or higher maps obtained from EBSD measurement of the specimens after the heat can treatments. Capdevila et al. [14] reported that the boundaries with misorientations of 15whole force Capdevila et al. [14] reported that the boundaries with misorientations of 15° or higher can force a a crack to change its propagation path during cleavage fracture to efficiently inhibit cleavage crack propagation. Therefore, the density high angle grainfracture boundaryto isefficiently an important consideration crack to change its propagation path of during cleavage inhibit cleavage tocrack control and characterize toughness. It is proposed that the mean equivalent diameter of the ferrite propagation. Therefore, the density of high angle grain boundary is an important consideration to effective grain with a tolerance angle of 15˝ ((MED)15˝ ) and the fraction of high angle boundary control and characterize toughness. It is proposed that the mean equivalent diameter of the ferrite dominate thewith density of high angle boundary. related data are summarized in Table 3. As shown effective grain a tolerance angle of 15° The ((MED) 15°) and the fraction of high angle boundary in Table 3, only the (MED)15˝ obtained under reheating at 1080 ˝ C was larger than that obtained under dominate the density of high angle boundary. The related data are summarized in Table 3. As reheating at 990 ˝ C or 1030 ˝ C. The fractions of high angle boundary in the specimens reheated at shown in Table 3, only the (MED)15° obtained under reheating at 1080 °C was larger than that 990 ˝ C and 1030 ˝ C were similar, whereas that of specimens reheated at 1080 ˝ C was significantly obtained under reheating at 990 °C or 1030 °C. The fractions of high angle boundary in the lesser. As the tempering temperature increased from 550 ˝ C to 680 ˝ C, the fraction of high angle specimens reheated at 990 °C and 1030 °C were similar, whereas that of specimens reheated at 1080 boundary changed slightly, but the low angle boundary fraction increased. Therefore, the impact °C was significantly lesser. As the of tempering from 550 energy decreased with the increase reheating temperature and temperingincreased temperatures (Table 2).°C to 680 °C, the fraction of high angle boundary changed slightly, but the low angle boundary fraction increased. Therefore, the impact energy decreased with the increase of reheating and tempering temperatures (Table 2). According to the study by Iza‐Mendia and Gutiérrez [6], the smallest microstructural unit size contributing to grain boundary strengthening was the effective grain with a tolerance angle of 2°. A small mean equivalent diameter of grain with a tolerance angle of 2° ((MED)2°) will result in a better strengthening effect. The (MED)2° under reheating at 1080 °C was larger than that under reheating
According to the study by Iza‐Mendia and Gutiérrez [6], the smallest microstructural unit size contributing to grain boundary strengthening was the effective grain with a tolerance angle of 2°. A small mean equivalent diameter of grain with a tolerance angle of 2° ((MED)2°) will result in a better strengthening effect. The (MED) 2° under reheating at 1080 °C was larger than that under reheating Metals 2016, 6, 75 9 of 11 at 990 or 1030 °C.
Figure figure maps obtained from EBSD measurement of theof specimens treatedtreated under Figure 8. 8. Inverse Inverse pole pole figure maps obtained from EBSD measurement the specimens ˝ C, TT = 550 ˝ C; (b) RT = 990 ˝ C, TT = 680 ˝ C; (c) RT = 1030 ˝ C, different conditions: (a) RT = 990 under different conditions: (a) RT = 990 °C, TT = 550 °C; (b) RT = 990 °C, TT = 680 °C; (c) RT = 1030 TT = 550 ˝ C; (d) RT = 1030 ˝ C, TT = 680 ˝ C; (e) RT = 1080 ˝ C; TT = 550 ˝ C; (f) RT = 1080 ˝ C, TT = 680 ˝ C °C, TT = 550 °C; (d) RT = 1030 °C, TT = 680 °C; (e) RT = 1080 °C; TT = 550 °C; (f) RT = 1080 °C, TT = (RT: reheating temperature; TT: tempering temperature). 680 °C (RT: reheating temperature; TT: tempering temperature). Table 3. EBSD results of the microstructural quantification. Table 3. EBSD results of the microstructural quantification. Reheating Tempering Fraction of Fraction of High Angle Fraction of Fraction of Low Angle Reheating Tempering Temperature Temperature Low Angle High Angle ˝ ˝ Boundaries, % Boundaries, % Temperature ( C) Temperature ( C) Boundaries, % Boundaries, % (°C) (°C) 550 15.8 57.3 550 15.8 57.3 990 990 680 21.4 58.7 680 21.4 58.7 550 11.2 58.8 550 11.2 58.8 1030 1030 680 16.3 54.6 680 16.3 54.6 550 17.2 36.3 1080 550 17.2 36.3 680 18.2 17.6 1080 680 18.2 17.6
(MED) (MED)2°2˝ (μm) (µm)
(MED) 15° ˝ (MED) 15 (μm) (µm)
2.7 2.7 2.8 2.8 2.3 2.3 2.5 2.5 3.9 3.9 4.1 4.1
3.5 3.5 3.3 3.3 3.2 3.2 3.1 3.1 4.8 4.8 4.7 4.7
In addition, in this work, we did not find large amounts of small alloyed precipitates with According to the study by Iza-Mendia and Gutiérrez [6], the smallest microstructural unit size dispersed distribution in all specimens. Moreover, the dislocation hardening was minor compared contributing to and grain boundary strengthening was thedislocation effective grain with a tolerance angleof ofthe 2˝ . to other terms during long‐term tempering, the density decreased because ˝ A small mean equivalent diameter of grain with a tolerance angle of 2 ((MED)2˝ ) will result in a better occurrence of recovery. Hence, the contribution of alloyed precipitation and dislocation density to strengthening effect. The (MED)2˝ under reheating at 1080 ˝ C was larger than that under reheating microhardness appeared to be similar for the specimens after different post‐weld heat treatments. ˝ C. at 990 or 1030 Therefore, grain size is a very important factor that affects the hardness. Figure 9 shows the In addition, in this work, we did not find large amounts of small alloyed precipitates with relationship of the values of microhardness with a function of MED at a tolerance angle of 2° to 15°. dispersed distribution in all specimens. Moreover, the dislocation hardening was minor compared to other terms and during long-term tempering, the dislocation density decreased because of the occurrence of recovery. Hence, the contribution of alloyed precipitation and dislocation density to microhardness appeared to be similar for the specimens after different post-weld heat treatments. Therefore, grain size is a very important factor that affects the hardness. Figure 9 shows the relationship
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The current results were in good agreement with this method, but when 8°, the fitting results of the values of microhardness with a function of MED at a tolerance angle of 2˝ to 15˝ . The current were less satisfying, which revealed that hardness was controlled ˝ by low misorientation grain results were in good agreement with this method, but when θ ě 8 , the fitting results were less boundaries. satisfying, which revealed that hardness was controlled by low misorientation grain boundaries.
Figure 9. Microhardness of the test steel varied with the reciprocal square root of effective grain size Figure 9. Microhardness of the test steel varied with the reciprocal square root of effective grain size with different tolerance angles: (a) 2˝ ; (b) 4˝ ; (c) 6˝ ; (d) 8˝ ; (e) 10˝ ; (f) 15˝ . with different tolerance angles: (a) 2°; (b) 4°; (c) 6°; (d) 8°; (e) 10°; (f) 15°.
4. Conclusions 4. Conclusions (1) At a higher reheating temperature, austenite grows rapidly and leads to a low-temperature (1) At a higher reheating temperature, austenite grows rapidly and leads to a low‐temperature microstructure. The low density of high angle grain boundary in the microstructure after tempering microstructure. The low density of high angle grain boundary in the microstructure after tempering deteriorates the impact toughness. deteriorates the impact toughness. (2) At higher (2) At higher reheating reheating temperature, temperature, aa large large amount amount of of M/A M/A constituents constituents are are obtained; obtained; they they develop into clusters and coarsen after tempering. An increased possibility of microcrack formation is develop into clusters and coarsen after tempering. An increased possibility of microcrack formation accompanied by a loss of impact toughness. is accompanied by a loss of impact toughness.
(3) Within the range of tempering temperature in the work, increasing the tempering temperature makes little difference to the low density of high angle grain boundary, but accelerates the decomposition of the M/A constituents.
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(3) Within the range of tempering temperature in the work, increasing the tempering temperature makes little difference to the low density of high angle grain boundary, but accelerates the decomposition of the M/A constituents. (4) EGS contributed most to hardness and had a good linear relationship with the microhardness in the range of tolerance angle 2˝ –8˝ . Acknowledgments: This work was financially supported by the Natural Science Foundation of China (51171162) and the Hebei Provincial Key Technology Support Program (14211010D, 15211019D). Author Contributions: The experiments were arranged by Xiu-Lin Han, Da-Yong Wu and Xiang-Ling Min. The experiments were conducted by Xu Wang, Xiu-Lin Han and Fu-Ren Xiao. The article was written by Xiu-Lin Han and revised by Bo Liao and Fu-Ren Xiao. Conflicts of Interest: The authors declare no conflict of interest.
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