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Mar 25, 2016 - 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, ...
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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