Metallurgical Effects of Niobium and Molybdenum on Heat-Affected Zone Toughness in Low-Carbon Steel

applied sciences

Article Metallurgical Eﬀects of Niobium and Molybdenum on Heat-Aﬀected Zone Toughness in Low-Carbon Steel

Hardy Mohrbacher 1,2 1 NiobelCon bvba, 2970 Schilde, Belgium 2 Department of Materials MTM, Leuven University (KU Leuven), 3001 Leuven, Belgium; [email protected]; Tel.: +32-3-484-5206 Received: 7 April 2019; Accepted: 29 April 2019; Published: 5 May 2019

Abstract: Modern weldable high strength steel grades are typically based on low-carbon alloy concepts using microalloying for obtaining a good strength-toughness balance. Such steel grades having a yield strength in the range of 420 to 690 MPa are very commonly used in pipelines, heavy vehicles, shipbuilding and general structural applications. Thermomechanical processing during hot rolling combined with accelerated cooling is an established means of producing such steel grades. Considering the alloying concepts, the use of niobium and molybdenum, and in selected cases boron, is very efficient to achieve high strength and good toughness. However, all targeted applications of such high strength steels involve extensive welding. Thus, heat affected zone properties are of particular importance. The present paper investigates the effects of Nb, Mo and Ti on the heat affected zone properties. Variations of the Mn and Si contents are considered as well. Additionally, the influence of post-weld heat treatment in the coarse-grained heat-affected zone (HAZ) is considered. In this approach, HAZ subzones were generated using laboratory weld cycle simulations in combination with systematic variation of alloying elements to scrutinize and interpret their specific effects. The results indicate that Mo and Nb, when alloyed in the typical range, provide excellent HAZ toughness and guarantee sufficiently low ductile-to-brittle transition temperature. An alloy combination of Nb, Mo and Ti improves performance under hot deformation conditions and toughness after post-weld heat treatment.

Keywords: welding simulation; heat-aﬀected zone; post-weld heat treatment; high-strength steel; microalloying; solubility; precipitation; solute drag; abnormal grain coarsening; hot ductility

1. Introduction The development of steel for use in pipeline, structural and automotive applications has seen a remarkable evolution since the late 1960s [1]. The increasing demands in terms of strength and toughness in general and weldability, in particular, have led to the design of alloys with reduced carbon content (Figure1a). Nowadays, steel of 700 MPa yield strength (grade 100) is achievable with a maximum carbon content of 0.05 percent [2]. Reducing carbon increases upper shelf energy and lowers the ductile-to-brittle transition temperature (Figure1b). The required strength is provided by either microstructural refinement, precipitation hardening or dislocation hardening. Typically a combination of these mechanisms is used. Since grain refinement is the only mechanism simultaneously improving strength and toughness (Figure1b), it should be always applied in the first place before employing other strengthening mechanisms. Microstructural refinement is efficiently achieved by niobium micro-alloying in combination with thermo-mechanical controlled processing (TMCP). This approach of producing high strength low alloy (HSLA) steel containing below 0.1 percent carbon and niobium micro-alloying between 0.02 and 0.04 percent is covering a wide portfolio of applications in the

Appl. Sci. 2019, 9, 1847; doi:10.3390/app9091847 www.mdpi.com/journal/applsci Appl. Sci. 2019, 9, 1847 2 of 17 Appl. Sci. 2019, 9, 1847 2 of 16 belowyield strength 0.1 percent range carbon of 300 and to 500 niobium MPa. Typically,micro-alloying the carbon between content 0.02is and between 0.04 percent 0.06 and is0.09 covering percent a widefor many portfolio of these of steelapplications grades resultingin the yield in a streng ferritic-pearliticth range of or 300 pearlite-reduced to 500 MPa. Typically, ferritic microstructure. the carbon contentThe trend is between of further 0.06 reducing and 0.09 the percent carbon for content many of is these reasoned steel bygrades improving resulting upper in a ferritic-pearlitic shelf toughness orand pearlite-reduced reducing ductile-to-brittle-transition ferritic microstructure. temperature The trend (DBTT). of further Besides, reducing avoiding the thecarbon peritectic content range is reasonedhas clear by beneﬁts improving with regardsupper shel to slabf toughness surface quality.and reducing In the ductile-to-brittle-transition 1980s, Hulka and Gray [3] temperature proposed an (DBTT).extraordinary Besides, alloy avoiding concept the using peritectic particularly range lowhas carbonclear benefits content with in combination regards to slab with surface a high niobiumquality. Inmicro-alloy the 1980s, addition Hulka and of 0.1 Gray percent. [3] proposed an extraordinary alloy concept using particularly low carbon content in combination with a high niobium micro-alloy addition of 0.1 percent.

(a) (b)

FigureFigure 1. DesignDesign criteria criteria for for weldable weldable high high strength strength steels: ( (aa)) The The Graville Graville diagram diagram classifying classifying weldabilityweldability of of steels steels and and historical evolution of of chemical composition composition represented represented by by the the carbon carbon equivalentequivalent (CE) (CE) in in high high strength strength steels; ( b)) ductile-to-brittle transition transition (DBTT) (DBTT) behavior behavior and and criteria criteria influencinginﬂuencing the the toughness toughness behavior behavior in in steel. steel.

TheThe intrinsic intrinsic benefits benefits of of modern low-carbon low-carbon alloy alloy concepts concepts are are a a particularly high high upper upper shelf shelf toughnesstoughness and and very very low low DBTT. DBTT. The The microstructure microstructure of of such such alloys alloys is is nearly nearly pearlite-free pearlite-free ferrite ferrite or or by by usingusing increased increased cooling rate, acicular ferritic. The The high high recrystallization recrystallization stop stop temperature temperature of of such such high-niobiumhigh-niobium alloyed alloyed steel steel principally principally widens widens the the processing processing window window for for TMCP TMCP rolling rolling and and enables enables higherhigher production production throughput, throughput, especially especially on on hot hot st striprip mills. mills. On On the the other other hand, hand, it it allows allows producing producing ferritic-bainiticferritic-bainitic microstructure microstructure even even by by air-cooling air-cooling after after finish finish rolling. rolling. Fine-tuning Fine-tuning of the of original the original low- carbonlow-carbon high-niobium high-niobium alloyed alloyed steel steel concept concept using using chromium chromium and and molybdenum molybdenum as as co-alloying elementselements has opened a widewide rangerange ofof interestinginteresting applications applications for for producing producing steel steel in in the the range range of of 460 460 to to700 700 MPa MPa yield yield strength. strength. These These can can be be found found in in pipe pipeas as wellwell asas automotive applications. Considering Considering structuralstructural applications, applications, some some traditional traditional specificatio specificationsns do do not not allow allow the the use use of of such such innovative innovative alloy alloy conceptsconcepts since since the the niobium niobium content content is is often often restricted restricted although although several several studies studies indicated indicated that that concerns concerns maymay not not be be justified justified [4–6]. [4–6]. Historically, Historically, the the limitati limitationon of of niobium addition addition re relateslates to to concerns concerns about about weldabilityweldability and and low low heat heat affected affected zone (HAZ) (HAZ) toughne toughness.ss. Thereby, Thereby, it it is is not not accounted accounted for for that that higher higher niobiumniobium additions additions are are typically used in extra-low-carbonextra-low-carbon steels. Previous investigations applying industrialindustrial weldingwelding processes processes to suchto such steels steels revealed revealed that heat-a that ffheat-affectedected zone toughness zone toughness is not negatively is not negativelyaffected by affected niobium by additions niobium inadditions the range in fromthe rang 0.02e from to 0.11% 0.02 [to7,8 0.11%]. The [7,8 present]. The study present will study focus will on focusthe metallurgical on the metallurgical effects and effects microstructural and microstructu changesral changes occurring occurring in the individualin the individual sub-zones sub-zones of the ofheat-a the ffheat-affectedected zone in zone extra-low-carbon in extra-low-carbon steel (0.035% steel (0.035% C) primarily C) primarily considering considering the effects the of effects niobium of niobiumand molybdenum and molybdenum alloying. alloying. The effects The of effects titanium of andtitanium silicon and alloying silicon are alloying additionally are additionally addressed. addressed.Post-weld heatPost-weld treatment heat (PWHT) treatment is used (PWHT) in selected is used applications in selected for weld applications stress relaxation for weld in thickerstress relaxation in thicker gaged components fabricated from high strength steel [9]. Micro-alloyed steels Appl. Sci. 2019, 9, 1847 3 of 17 Appl. Sci. 2019, 9, 1847 3 of 16 gaged components fabricated from high strength steel [9]. Micro-alloyed steels can show embrittlement can show embrittlement by such treatment. Thus, simulations of PWHT are performed in this study by such treatment. Thus, simulations of PWHT are performed in this study to investigate the to investigate the embrittlement sensitivity of these steels in the coarse-grained heat affected zone. embrittlement sensitivity of these steels in the coarse-grained heat affected zone. 2. Materials and Methods 2. Materials and Methods

2.1.2.1. Definition Deﬁnition of of the the Test Test Materials Low-carbonLow-carbon (0.035%) (0.035%) steels steels with with a a systematic systematic variation variation of of alloyi alloyingng elements elements were were designed designed as as specifiedspeciﬁed in in Table Table 11.. SuchSuch alloysalloys areare beingbeing usedused forfor pipepipe steelssteels withwith highhigh requirementsrequirements regardingregarding toughnesstoughness and and sour sour gas gas resistance. resistance. Several Several of of these these al alloyloy compositions compositions are are in in actual actual industrial industrial use use for for pipepipe related related strip strip and and plate plate products. products. Thereby, Thereby, lower-Mn lower-Mn and and Mo-free Mo-free variants variants cover cover X60–X65 X60–X65 grades grades whereas thethe higherhigher alloyedalloyed variants variants are are being being used used for for X70 X70 and and X80 X80 grades grades (X refers (X refers to minimum to minimum yield yieldstrength, strength, number number in ksi). in ksi).

TableTable 1. VariationVariation of alloy compositions in mass mass % % subjected subjected to to heat-affected heat-aﬀected zone zone (HAZ) (HAZ) simulations simulations (symbols(symbols refer refer to to all all subsequent diagrams).

Nb Alloy Symbol C Si Mn Cr Ti Mo (varied in all alloys)

V1 0.00 1.45 V2 0.00 a: 0.00 V3 0.30 0.015 b: 0.02 0.035 0.20 c: 0.05 V4 d: 0.10 1.90 V5 0.00 0.15

V6 0.10 0.015

All alloy variants (V1–V6) were produced as 300 kg laboratory melts in an induction furnace. Each Allsub-variant alloy variants (a–d) (V1–V6)was cast were into produceda 75 kg ingot as 300 and kg laboratorythe niobium melts level in was an induction increased furnace. in the remainingEach sub-variant melt to (a–d) the wasnext castlevel. into The a 75 ingots kg ingot were and forged the niobium into square level bars was increased(75 × 75 mm) in the at remaining 1200 °C. melt to the next level. The ingots were forged into square bars (75 75 mm) at 1200 C. These bars These bars were subsequently cut into segments, which were subjected× to the following◦ rolling schedule:were subsequently cut into segments, which were subjected to the following rolling schedule: • Reheating temperature: 1150 °C +/−10 °C. Reheating temperature: 1150 ◦C +/ 10 ◦C. • • Rouging rolling in four passes− to 35 mm. Rouging rolling in four passes to 35 mm. • • Thermo-mechanical rolling to 13 mm in five passes. Thermo-mechanical rolling to 13 mm in five passes. • • Final rolling temperature: 850 °C +/−20 °C Final rolling temperature: 850 C +/ 20 C • • Water cooling or in selected◦ − case◦ air cooling. Water cooling or in selected case air cooling. • The produced steel plates were characterized by tensile testing and the resulting values are displayedThe producedin Figure steel2a. The plates alloy were variants characterized cover a yield by tensilestrength testing (Rp0.2) andrange the from resulting 350 tovalues 600 MPa. are Thereby the increase of the Nb level in each variant clearly relates to a strength increase, amounting displayed in Figure2a. The alloy variants cover a yield strength (R p0.2) range from 350 to 600 MPa. toThereby 50–160 the MPa increase when ofadding the Nb 0.1%Nb. level in eachThe tensile variant st clearlyrengthrelates correlates to a strengthwell with increase, the yield amounting strength indicatingto 50–160 MPaan average when addingyield-to-tensile 0.1%Nb. ratio The (YTR) tensile of strength 0.78. The correlates toughness well of witheach thealloy yield variant strength was characterizedindicating an by average standard yield-to-tensile Charpy V-notch ratio testing (YTR) using of 0.78. a 10 The × 10 toughness mm sample of size each (acc. alloy to DIN variant 10045). was Thecharacterized upper-shelf by toughness standard Charpy of these V-notch steel alloys testing scatte usingrs aaround 10 10 200 mm J. sampleThis level size is (acc. relatively to DIN low 10045). for × suchTheupper-shelf low-carbon toughness steel. An of theseindustrially steel alloys produced scatters aroundreference 200 material J. This level (X65) is relatively having lownearly for suchthe compositionlow-carbon steel.of variant An industrially V2-c showed produced an upper-shel referencef toughness material (X65) of 370 having J in comparative nearly the composition tests. The ductile-to-brittleof variant V2-c showed transition an temperature, upper-shelf toughness i.e., the temperature of 370 J in comparativeat which 50% tests. of the The upper-shelf ductile-to-brittle energy istransition reached temperature,(DBTT T50%, i.e., Figure the temperature1b) is also lower at which in the 50% industrial of the upper-shelf material (–125 energy °C) is as reached compared (DBTT to the laboratory alloy V2-c (–85 °C). The differences in the toughness behavior are due to a lower T50%, Figure1b) is also lower in the industrial material ( 125 ◦C) as compared to the laboratory alloy cleanness of the laboratory melts, which were produced− under non-vacuum conditions. However, V2-c ( 85 ◦C). The differences in the toughness behavior are due to a lower cleanness of the laboratory the comparison− between the various laboratory melts allows identifying the toughness influence of the alloying elements relative to each other. The DBTT results of all laboratory steels are shown in Figure 2b. It is apparent that the DBTT in each alloy system decreases with increasing Nb addition, Appl. Sci. 2019, 9, 1847 4 of 17 melts, which were produced under non-vacuum conditions. However, the comparison between the various laboratory melts allows identifying the toughness influence of the alloying elements relative to Appl.each Sci. other. 2019, The9, 1847 DBTT results of all laboratory steels are shown in Figure2b. It is apparent that4 of the 16 DBTT in each alloy system decreases with increasing Nb addition, irrespective of the strength level of irrespectivethe base alloy. of This the beneficial strength behaviorlevel of isthe related base to al theloy. pronounced This beneficial microstructural behavior refinement is related induced to the pronouncedby Nb micro-alloying. microstructural The tworefine datament points induced marked by Nb as micro-alloying. “PF” in Figure 2Theb have two adata mostly points polygonal marked asferritic “PF” (PF) in Figure microstructure 2b have a in mostly contrast polygonal to the acicular ferritic ferritic (PF) microstructure microstructure in of contrast all other to alloys. the acicular ferritic microstructure of all other alloys.

(a) (b)

Figure 2. Properties of the as-rolled materials: ( a) Relationship of yieldyield strengthstrength (R(Rp0.2p0.2)) and and tensile strength;strength; (b(b)) ductile-to-brittle ductile-to-brittle transition transition temperature temperature (50%) (50%) determined determined by Charpy by testingCharpy (PF: testing polygonal (PF: polygonalferrite matrix, ferrite all matrix, other acicular all other ferrite acicular matrices). ferrite matrices).

The base material material microstructure microstructure and the refining refining effect effect of Nb are demonstrated by light optical optical micrographs in Figure 33.. The Nb-freeNb-free basebase alloyalloy V2-aV2-a revealsreveals aa polygonalpolygonal ferriticferritic microstructuremicrostructure (Figure 33a)a) containingcontaining carbon-richcarbon-rich phasesphases suchsuch asas degenerateddegenerated pearlitepearlite (dark(dark contrast)contrast) andand aa smallsmall amount ofof MA-phase MA-phase (white (white islands). islands). The microstructureThe microstructure of the Nb-alloyedof the Nb-alloyed variant (V2-d) variant is dominated(V2-d) is dominatedby fine-grained by fine-grained acicular ferrite. acicular The ferrite. addition The ofaddition molybdenum of molybdenum to Nb-free to steelNb-free (V4-a) steel prevents (V4-a) preventsthe formation the formation of polygonal of ferritepolygonal and promotesferrite and a coarse-grainedpromotes a coarse-grained acicular ferritic acicular microstructure ferritic microstructure(Figure3c). Adding (Figure niobium 3c). Adding to this niobium alloy (V4-d) to this results alloy (V4-d) in remarkable results in refinement remarkable of refinement this acicular of thisferritic acicular microstructure ferritic microstructure (Figure3d). (Figure 3d).

Appl. Sci. 2019, 9, 1847 5 of 17 Appl. Sci. 2019, 9, 1847 5 of 16 Appl. Sci. 2019, 9, 1847 5 of 16

(a) (b) (a) USV Symbol(b) Macro(s) Description

2934 ⤴ \textrcurvearrowne ARROW POINTING RIGHTWARDS THEN CURVING UPWARDS 2935 ⤵ \textlcurvearrowse ARROW POINTING RIGHTWARDS THEN CURVING DOWNWARDS 2936 ⤶ \textlcurvearrowsw ARROW POINTING DOWNWARDS THEN CURVING LEFTWARDS 2937 ⤷ \textrcurvearrowse ARROW POINTING DOWNWARDS THEN CURVING RIGHTWARDS 2938 ⤸ \textlcurvearrowdown RIGHT-SIDE ARC CLOCKWISE ARROW 2939 ⤹ \textrcurvearrowdown LEFT-SIDE ARC ANTICLOCKWISE ARROW 293A ⤺ \textrcurvearrowleft TOP ARC ANTICLOCKWISE ARROW 293B ⤻ \textrcurvearrowright BOTTOM ARC ANTICLOCKWISE ARROW (c) 294A ⥊ (d) \textleftrightharpoon LEFT BARB UP RIGHT BARB DOWN HARPOON (c) 294B ⥋ (d) \textrightleftharpoon LEFT BARB DOWN RIGHT BARB UP HARPOON Figure 3. Light optical micrographs (Le Pera etching) of the294C as-rolled⥌ materials:\textupdownharpoonrightleft (a) V2-a (polygonal UP BARB RIGHT DOWN BARB LEFT HARPOON Figure 3. Light optical micrographs (Le Pera etching)etching) of the as-rolledas-rolled materials:materials: (a) V2-a (polygonal ferrite); (b) V2-d (mostly acicular ferrite); (c) V4-a (coarse294D acicular ⥍ferrite);\textupdownharpoonleftright (d) V4-d (fine acicular UP BARB LEFT DOWN BARB RIGHT HARPOON ferrite); ((bb)) V2-dV2-d (mostly (mostly acicular acicular ferrite); ferrite); (c) V4-a(c) V4-a (coarse (coarse acicular acicular ferrite); ferrite); (d) V4-d (d) (fine V4-d acicular (fine acicular ferrite). ferrite). 2962 ⥢ \textleftleftharpoons LEFTWARDS HARPOON WITH BARB UP ABOVE LEFTWARDS HARPOON WITH BARB ferrite). DOWN 2963 ⥣ \textupupharpoons UPWARDS HARPOON WITH BARB LEFT BESIDE UPWARDS HARPOON WITH BARB 2.2. Heat Affected Zone Simulation and Characterization RIGHT 2.2. Heat Affected Zone Simulation and Characterization 2964 ⥤ \textrightrightharpoons RIGHTWARDS HARPOON WITH BARB UP ABOVE RIGHTWARDS HARPOON WITH BARB 2.2. HeatThe Affected simulation Zone of Si heatmulation affected and Characterization sub-zones was performed in a HAZ simulator (Vatron GmbH, DOWN The simulation of heat affected sub-zones was performed2965 in a⥥ HAZ \textdowndownharpoonssimulator (Vatron GmbH, DOWNWARDS HARPOON WITH BARB LEFT BESIDE DOWNWARDS HARPOON WITH Linz,The Austria) simulation using electricof heat currentaffected for sub-zones fast conductive was performed heating to in a peaka HAZ temperature simulator and(Vatron an air-water GmbH, BARB RIGHT 296A ⥪ \textleftbarharpoon LEFTWARDS HARPOON WITH BARB UP ABOVE LONG DASH mixtureLinz, Austria) applied using via four electric coaxially current positioned for fast nozzles conductive for rapid heating cooling to a (Figure peak 4temperature). A thermocouple and an fixed air- Linz, Austria) using electric current for fast conductive heating296B to⥫ a peak\textbarleftharpoon temperature and an air- LEFTWARDS HARPOON WITH BARB DOWN BELOW LONG DASH inwater the centermixture of theapplied sample via provides four coaxially feedback posi for thetioned time-temperature nozzles for rapid control cooling software (Figure (Lab-View). 4). A water mixture applied via four coaxially positioned nozzles for rapid cooling (Figure 4). A RIGHTWARDS HARPOON WITH BARB UP ABOVE LONG DASH thermocouple fixed in the center of the sample provides 296Cfeedback⥬ for the\textrightbarharpoon time-temperature control Thethermocouple samples for fixed this testin the were center machined of the fromsample the provides laboratory-rolled feedback296D 13 ⥭for mm the thick\textbarrightharpoon time-temperature plates to a size control of 10 RIGHTWARDS HARPOON WITH BARB DOWN BELOW LONG DASH software (Lab-View).3 The samples for this test were machined from the laboratory-rolled 13 mm thick× 10software95 mm (Lab-View).. These samples The samples were instrumentedfor this test were with machined thermocouples296E from the⥮ and laboratory-rolled clamped\textupdownharpoons into the 13 electrodesmm thick UPWARDS HARPOON WITH BARB LEFT BESIDE DOWNWARDS HARPOON WITH BARB plates× to a size of 10 × 10 × 95 mm3. These samples were instrumented with thermocouples and RIGHT platesleaving to a a free size length of 10 of× 6010 mm.× 95 mm This3. setupThese results samples in awere sample296F instrumented volume⥯ of\textdownupharpoons with approximately thermocouples 10 10and DOWNWARDS HARPOON WITH BARB LEFT BESIDE UPWARDS HARPOON WITH BARB × × RIGHT 15clamped mm3 having into the a electrodes homogeneous leaving heat a afreeffected length microstructure. of 60 mm. This For setup hardness results measurements in a sample volume (Vickers of clamped into the electrodes leaving a free length of 60 mm.2987 This setup⦇ results\textllparenthesis in a sample volume of Z NOTATION LEFT IMAGE BRACKET approximately 10 × 10 × 15 mm3 having a homogeneous heat affected microstructure. For hardness approximatelyHV10) and metallographic 10 × 10 × 15 mm analysis,3 having the a homogeneous homogeneous HAZ heat2988 wasaffected cut⦈ frommicrostructure.\textrrparenthesis the sample. For For hardness tensile Z NOTATION RIGHT IMAGE BRACKET measurements (Vickers HV10) and metallographic analysis, the homogeneous HAZ was cut from the measurementstesting, the bar (Vickers was machined HV10) toand a roundmetallographic standard analysis, tensile29B0 sample the homogeneous (⦰ 10 mm)\textinvdiameter with HAZ the was homogeneous cut from the REVERSED EMPTY SET sample. For tensile testing, the bar was machined to a round standard tensile sample2 (⦰ 10 mm) with zone being centered in the gage length. Likewise, Charpy29B6 test bars⦶ of 10 \textobar10 mm were machined CIRCLED VERTICAL BAR sample. For tensile testing, the bar was machined to a round standard tensile sample (⦰ 10 mm) with2 the homogeneous zone being centered in the gage length.29B8 Likewise,⦸ Charpy\textobslash× test bars of 10 × 10 mm CIRCLED REVERSE SOLIDUS thehaving homogeneous the V-notch zone centered being to centered the homogeneous in the gage HAZ. length. Likewise, Charpy test bars of 10 × 10 mm2 were machined having the V-notch centered to the homogeneous29BA HAZ.⦺ \textobot CIRCLE DIVIDED BY HORIZONTAL BAR AND TOP HALF DIVIDED BY VERTICAL BAR were machined having the V-notch centered to the homogeneous HAZ. 29BB ⦻ \textNoChemicalCleaning CIRCLE WITH SUPERIMPOSED X 29C0 ⧀ \textolessthan CIRCLED LESS-THAN 29C1 ⧁ \textogreaterthan CIRCLED GREATER-THAN

29C4 ⧄ \textboxslash SQUARED RISING DIAGONAL SLASH 29C5 ⧅ \textboxbslash SQUARED FALLING DIAGONAL SLASH

29C6 ⧆ \textboxast SQUARED ASTERISK 29C7 ⧇ \textboxcircle SQUARED SMALL CIRCLE

29C8 ⧈ \textboxbox SQUARED SQUARE 29D3 ⧓ \textValve BLACK BOWTIE 29DF ⧟ \textmultimapboth DOUBLE-ENDED MULTIMAP 29E2 ⧢ \textshuffle SHUFFLE PRODUCT

2A04 ⨄ \textuplus N-ARY UNION OPERATOR WITH PLUS TWO LOGICAL AND OPERATOR (a) 2A07 ⨇(b) \textbigdoublewedge (a) 2A08 ⨈(b) \textbigdoublevee TWO LOGICAL OR OPERATOR Figure 4. (a)(a) Setup for HAZ simulation and ((b)b) sample 2A1D geometries ⨝ for mechanical\textJoin testing. JOIN Figure 4. (a) Setup for HAZ simulation and (b) sample geometries for mechanical testing. 2A1F ⨟ \textfatsemi Z NOTATION SCHEMA COMPOSITION HAZ simulations were performed for the inter-critical2A22 zone ⨢(ICZ), \textcircplusi.e., the peak temperature PLUS SIGN WITH SMALL CIRCLE ABOVE HAZ simulations were performed for the inter-critical zone (ICZ), i.e., the peak temperature being in the ferrite-austenite two-phase region, and for the2A2A coarse-grained⨪ \textminusdot zone (CGZ) using a peak MINUS SIGN WITH DOT BELOW being in the ferrite-austenite two-phase region, and for the2A30 coarse-grained⨰ \textdottimes zone (CGZ) using a peak MULTIPLICATION SIGN WITH DOT ABOVE 2A32 ⨲ \textdtimes SEMIDIRECT PRODUCT WITH BOTTOM CLOSED 2A38 ⨸ \textodiv CIRCLED DIVISION SIGN 2A3C ⨼ \textinvneg INTERIOR PRODUCT 2A40 ⩀ \textcapdot INTERSECTION WITH DOT 2A4E ⩎ \textsqdoublecap DOUBLE SQUARE INTERSECTION 2A4F ⩏ \textsqdoublecup DOUBLE SQUARE UNION

48 Appl. Sci. 2019, 9, 1847 6 of 17

HAZ simulations were performed for the inter-critical zone (ICZ), i.e., the peak temperature being in the ferrite-austenite two-phase region, and for the coarse-grained zone (CGZ) using a peak temperature of 1300 ◦C. For defining the peak temperature in the ICZ, dilatometer tests were done identifying the transformation points in the up-heat cycle at a heating rate of 600 K/s. The peak temperature should be set to in between the maximum transformation points Ac1p and Ac3p. This condition was for all alloys fulfilled for a peak temperature of 900 ◦C. The fine-grained zone requires a peak temperature being approximately 50 ◦C above the transformation finish temperature Ac3e. This condition was fulfilled at 1100 ◦C for the current alloys. For obtaining the characteristic tensile properties across the entire HAZ, samples were produced at peak temperatures of 300, 600, 800, 900, 1100 and 1300 ◦C and machined into tensile bars. The heating rate was always set to around 600 K/s and the cooling rate was adjusted to yield an 800–500 ◦C temperature interval (∆T8/5) of 25 s corresponding to typical values obtained by submerged arc welding in practice.

2.3. Post-Weld Heat Treatment Simulation and Characterization For simulating the material behavior under post-weld heat treatment (PWHT), the coarse-grained microstructure was produced in a ﬁrst step using a peak temperature of 1300 ◦C as described above. Machined round tensile bars from this CGZ material were then subjected to hot tensile tests in the temperature range of 300 to 620 ◦C. Particularly the reduction of area (RA) measured under the given tensile temperature is taken as an indicator for relaxation embrittlement sensitivity. In industrial applications, PWHT is typically performed in the temperature range of 500 to 600 ◦C[9]. To characterize the material behavior after such a treatment in detail, CGZ samples were heat treated at 560 ◦C for 30 min. Thereafter samples of all alloy variants were characterized in the same fashion as described for the HAZ simulations providing hardness, tensile test and toughness data.

3. Results The experimental program as described in paragraph 2 delivered a large amount of data of which only part is reported here allowing to draw relevant conclusions on the eﬀect of alloying elements such as niobium, molybdenum, manganese and silicon. Regarding toughness, the upper-shelf energy values measured in the ICZ and CGZ are similar to those found in the base material. Again, the industrial reference steel X65 shows in comparison to laboratory material V2-c much higher upper-shelf energies and lower DBTT values in the ICZ and CGZ. A thorough analysis of all toughness data revealed that the DBTT T50% data exhibit the strongest correlation with the alloy variations. Therefore, only DBTT T50% data are considered in the following analysis.

3.1. Mechanical Characteristics of HAZ Sub-Zones Figure5 indicates the change in mechanical characteristics for alloy V3 after exposure to peak temperatures in the range of 300 to 1300 ◦C. The hardness measurements indicate signiﬁcant softening in the ICZ, i.e., after heating to 900 ◦C peak temperature. The softening is most pronounced for the Nb-free steel losing 72 HV10 compared to the base material. For the alloy with the highest Nb addition, the hardness drop is only 36 HV10. In the CGZ, the behavior is similar yet less strong. Here the Nb-free alloy is softening by 52 HV10 while the highest Nb-alloyed variant shows no hardness loss. Already the smallest Nb addition of 0.02% reveals a signiﬁcant softening resistance in both, ICZ and CGZ. The measured HV10 hardness relates quite well to the measured tensile strength by a factor of 3.2. Instead of reporting yield and tensile strength data as a function of the peak temperature, it is more interesting looking at the evolution of the yield-to-tensile ratio (YTR). The latter is more conclusive with regard to microstructural changes. In Figure5 the YTR evolution is plotted as an interpolation of the data obtained at 300, 600, 800, 900, 1100 and 1300 ◦C for the Nb-free and 0.1%Nb containing alloy V3. Generally, the YTR of the 0.1%Nb steel is consistently higher than that of the Nb-free steel. The YTR increases form its base material when the peak temperature is raised to sub-critical values (

appearsAppl. Sci. at 2019 a somewhat, 9, 1847 higher temperature in the 0.1%Nb alloy as compared to the Nb-free alloy. In7 theof 16 latter, the YTR shows a sharp drop to a pronounced minimum value of 0.54 in the ICZ. The YTR of the 0.1%NbYTR of alloy,the 0.1%Nb on the alloy, contrary, on the drops contrary, only todrops 0.75 only in the to ICZ.0.75 in In the the ICZ. CGZ, In the the YTR CGZ, increases the YTR fromincreases its minimumfrom its minimum value in the value ICZ in to the a level ICZ betweento a level 0.7 between and 0.8. 0.7 and 0.8. TheThe increase increase of of the the YTR YTR in in the the sub-critical sub-critical temperature temperature range range of 600–800of 600–800◦C is°C typically is typically related related to temperingto tempering and precipitationand precipitation phenomena. phenomena. Tempering Tempering of hard of carbon-enriched hard carbon-enriched phases in phases the original in the materialoriginal reduces material primarily reduces theprimarily tensile the strength tensile while strength precipitation while precipitation of solute micro-alloys of solute micro-alloys primarily increasesprimarily the increases yield strength. the yield Particularly strength. theParticularly precipitation the precipitation behavior of Nb behavior will be discussed of Nb will in be paragraph discussed 4in moreparagraph detail. 4 Thein more very detail. low YTR The of very the Nb-free low YTR steel of the in the Nb-free ICZ is steel a clear in indicatorthe ICZ is for a clear a microstructure indicator for witha microstructure dual-phase character. with dual-phase The dispersion character. of The the martensiticdispersion of phase the martensitic in a soft matrix phase provides in a soft amatrix low yieldprovides and higha low tensile yield and strength. high tensile Such dual-phase strength. Su steelsch dual-phase are deliberately steels are produced deliberately by annealing produced in by theannealing inter-critical in the temperature inter-critical range temperature followed range by fast followed cooling. by Metallographic fast cooling. Metallographic analysis using Le analysis Pera etchantusing revealsLe Pera the etchant presence reveals of martensite the presence in the ICZof material.martensite In in the the optical ICZ microscope, material. In these the phasesoptical appearmicroscope, with a these white phases or brownish appear color. with a white or brownish color.

(c)

(a) (b)

FigureFigure 5. 5.Heat-a Heat-affectedﬀected zone zone simulation simulation of of alloy alloy V3: V3: (a ()a Evolution) Evolution of of hardness hardness and and yield-to-tensile yield-to-tensile ratioratio (YTR) (YTR) and and measured measured transformation transformation temperatures temperatures during during up-heating; up-heating; (b ()b microstructure) microstructure in in the the inter-criticalinter-critical zone zone (ICZ) (ICZ) of of the the Nb-free Nb-free alloy alloy V3-a; V3-a; (c ()c microstructure) microstructure of of alloy alloy V3-d V3-d containing containing 0.10%Nb 0.10%Nb (etching(etching with with Le Le Pera Pera agent). agent).

FigureFigure5b 5b shows shows the the presence presence of of a smalla small fraction fraction of of martensite martensite islands islands dispersed dispersed in in a fullya fully recrystallizedrecrystallized polygonal polygonal ferrite ferrite matrix matrix of the Nb-freeof the steel.Nb-free The originalsteel. The acicular original ferritic acicular microstructure ferritic ofmicrostructure the base material of the has base completely material has disappeared. completely In disappeared. the 0.1%Nb-alloyed In the 0.1%Nb-alloyed steel (Figure5 c)steel the (Figure ICZ microstructure5c) the ICZ microstructure also appears fully also recrystallized.appears fully However,recrystallized. the polygonal However, ferrite the polygonal grains are ferrite much grains finer thanare much those finer in the than Nb-free those alloy.in the Nb-free The amount alloy. of The martensite amount of islands martensite in the islands microstructure in the microstructure is almost negligible.is almost Grainnegligible. refinement Grain hasrefinement a stronger has eff aect stronger (Hall-Petch effect relationship) (Hall-Petchon relationship) the yield strength on the thanyield onstrength the tensile than strength.on the tensile Thus, strength. with decreasing Thus, with graindecreasing size the grain YTR size increases. the YTR increases. The strength The strength of the refinedof the polygonalrefined polygonal ferrite is ferrite however is however still lower still compared lower compared to the acicular to the acicular ferritic baseferritic material base material as the sub-structureas the sub-structure within the within grain the is lostgrain due is lost to recrystallization. due to recrystallization. AddingAdding molybdenum molybdenum in in alloy alloy V4 V4 shows shows qualitatively qualitatively the the same same mechanical mechanical behavior behavior as as in in alloy alloy V3V3 (Figure (Figure6a) 6a) and and similar similar e ff ectseffects of niobium.of niobium. However, However, the hardnessthe hardness scattering scattering range range between between Nb-free Nb- andfree 0.1%Nb and 0.1%Nb containing containing alloys inalloys the ICZin the and ICZ CGZ and is lessCGZ wide. is less Particularly wide. Particularly the Nb-free the alloyNb-free softens alloy softens less, which is a known effect caused by Mo. The 0.1%Nb steel, on the contrary, shows a greater hardness loss in the CGZ compared to the base material. This can indicate less precipitation hardening by niobium in the CGZ. The YTR indicates less softening in the sub-critical zone of the Nb- Appl. Sci. 2019, 9, 1847 8 of 17 less, which is a known effect caused by Mo. The 0.1%Nb steel, on the contrary, shows a greater hardness loss in the CGZ compared to the base material. This can indicate less precipitation hardening by niobium in the CGZ. The YTR indicates less softening in the sub-critical zone of the Nb-free alloy while the dual phase character in the ICZ (YTR = 0.54) is identical to alloy V3. The YTR in the CGZ is Appl.somewhat Sci. 2019, lower 9, 1847 as in alloy V3, regardless of the Nb content. 8 of 16 Reducing the silicon content in alloy V6 leads to somewhat more softening of the Nb-free alloy in the ICZReducing as compared the silicon to alloy content V4. in The alloy hardness V6 leads is howeverto somewhat still highermore softening than in alloy of the V3 Nb-free for the alloy same incondition. the ICZ as The compared 0.1%Nb to containing alloy V4. The steel hardness V6 shows is however no softening still higher in the than CGZ. in Regarding alloy V3 for the the YTR, same a condition.comparably The higher 0.1%Nb minimum containing value steel of 0.57V6 shows is observed no softening for the Nb-freein the CGZ. alloy Regarding in the ICZ the indicating YTR, a comparablya less pronounced higher minimum dual phase value character. of 0.57 Foris observed the 0.1%Nb for the containing Nb-free alloy steel, in the the YTR ICZin indicating the ICZ isa lessalso pronounced higher than dual in the phase previous charac alloys.ter. For Although the 0.1%Nb these containing differences steel, appear the toYTR be in relatively the ICZ modest, is also higherthe observed than in e fftheects previous can be attributedalloys. Although to the siliconthese differences alloy variation. appear In to dual be phaserelatively steel modest, production the observedsilicon is effects standardly can be used attributed to promote to the the silicon partitioning alloy variation. of carbon In dual to inter-critical phase steel austenite production and, silicon thus, isto standardly stabilize it used for transformation to promote the into partitioning martensite of upon carbon quenching. to inter-critical This e ffausteniteect is based and, on thus, silicon to stabilizeincreasing it thefor activitytransformation of carbon into and alsomartensite preventing upon the quenching. precipitation This of cementite.effect is based The low-Si on silicon alloys increasingexhibit the the highest activity YTR of valuescarbon in and the also CGZ. preventing the precipitation of cementite. The low-Si alloys exhibit the highest YTR values in the CGZ. The transformation temperature Ac1P is very similar for the three Nb-free variants of V3, V4 and The transformation temperature Ac1P is very similar for the three Nb-free variants of V3, V4 and V6 ranging between 804 and 810 ◦C whereas the A3cp temperature is lower for the Mo-alloyed variants. V6 ranging between 804 and 810 °C whereas the A3cp temperature is lower for the Mo-alloyed The addition of 0.1%Nb clearly shifts the Ac1p temperature upwards, particularly in combination with variants. The addition of 0.1%Nb clearly shifts the Ac1p temperature upwards, particularly in Mo-alloying. This is also the case for the Ac3p temperature in alloys V3-d and V4-d. Remarkably, combination with Mo-alloying. This is also the case for the Ac3p temperature in alloys V3-d and V4-d. the combined alloying of Mo and 0.1%Nb at reduced Si content leads to a much lower Ac3p temperature, Remarkably,diminishing the combined transition rangealloying to of only Mo 37 and K while0.1%Nb it is at between reduced 110Si content and 130 leads K for to the a much other lower alloys A(Figuresc3p temperature,5 and6). diminishing the transition range to only 37 K while it is between 110 and 130 K for the other alloys (Figures 5 and 6).

(a) (b)

FigureFigure 6. 6. EvolutionEvolution of of hardness hardness and and yi yield-to-tensileeld-to-tensile ratio ratio (YTR) (YTR) in in the the heat-affected heat-aﬀected zone zone of of ( (aa)) alloys alloys V4;V4; ( (bb)) alloys alloys V6. V6. Measured Measured transformation transformation temper temperaturesatures during during up-hea up-heatingting are are indicated. indicated.

3.2.3.2. Toughness Toughness Characteristics Characteristics of of HAZ HAZ Sub-Zones Sub-Zones Principally,Principally, the the transition transition from from ductile ductile to tobrittle brittle fracture fracture mode mode occurs occurs in bcc-type in bcc-type steel steel when when the yieldthe yield strength strength approaches approaches the cleavage the cleavage stress stress level level [10]. [Lower10]. Lower temperature temperature and higher and higher strain strain rate increaserate increase the yield the yield strength strength so that so under that under impact impact loading loading conditions conditions at a certain at a certain low lowtemperature temperature the yieldthe yield strength strength exceeds exceeds the thecleavage cleavage stress stress causing causing brittle brittle failure failure as as a aconsequence. consequence. Refining Refining the the microstructure also increases the yield strength (Hall-Petch relationship), yet the refinement has an even larger effect in terms of increasing the cleavage stress. Hence, microstructural refinement effectively lowers the DBTT. The yield strength is also increased by solid solution, precipitation and dislocation strengthening. These mechanisms have an adverse effect on the DBTT. Accordingly, for evaluating the alloy influence on the heat affected zone toughness it is most interesting to analyze the transition temperature (DBTT T50%) as a function of the actual yield strength in the inter-critical and coarse-grained sub-zones. Appl. Sci. 2019, 9, 1847 9 of 17 microstructure also increases the yield strength (Hall-Petch relationship), yet the refinement has an even larger effect in terms of increasing the cleavage stress. Hence, microstructural refinement effectively lowers the DBTT. The yield strength is also increased by solid solution, precipitation and dislocation strengthening. These mechanisms have an adverse effect on the DBTT. Accordingly, for evaluating the alloy influence on the heat affected zone toughness it is most interesting to analyze the transition temperature (DBTT T50%) as a function of the actual yield strength in the inter-critical Appl.and Sci. coarse-grained 2019, 9, 1847 sub-zones. 9 of 16 Figure7a displays the measured DBTT values as a function of the yield strength for the ICZ. The mostFigure obvious 7a displays observation the measured is that theDBTT DBTT values decreases as a function with increasing of the yield Nb strength content, for despite the ICZ. strongly The mostincreasing obvious yield observation strength. This is that can bethe due DBTT to pronounced decreases with grain increasing refinement Nb by Nbcontent, in the despite ICZ (Figure strongly5b,c). increasingMolybdenum yield alloyed strength. variants This tendcan be to due have to a pron higherounced DBTT grain at a given refinement yield strength by Nb in level the thanICZ Mo-free(Figure 5b,c).alloys. Molybdenum This discrepancy, alloyed however, variants diminishes tend to have with a increasinghigher DBTT Nb content.at a given The yield Mo-alloyed strength variantslevel than at Mo-freelower Nb alloys. contents This contain discrepancy, an increased however, amount diminish of carbon-riches with increasing phase, Nb which content. probably The isMo-alloyed tempered variantsmartensite. at lower The highest Nb contents DBTT contain is observed an increase for the steeld amount without of carbon-rich micro-alloy (nophase, Ti, nowhich Nb). probably The low-Si is temperedMo-alloyed martensite. variant, however, The highest shows DBTT a low is observed DBTT in thefor micro-alloy-freethe steel without state. micro-alloy (no Ti, no Nb). The low-Si Mo-alloyed variant, however, shows a low DBTT in the micro-alloy-free state.

(a) (b)

FigureFigure 7. 7. Ductile-to-brittleDuctile-to-brittle transition transition temperature temperature (DB (DBTTTT T50%) T50%) determined determined by Charpy by Charpy testing: testing: (a)

Simulated(a) Simulated intercritcal intercritcal HAZ HAZ (ICZ) (ICZ) wi withth a peak a peak temperature temperature of of 900 900 °C;◦C; (b (b) )simulated simulated coarse-grained coarse-grained HAZHAZ (CGZ) (CGZ) with with a a peak peak temperature temperature of 1300 °C.◦C.

InIn the the CGZ CGZ the the DBTT DBTT values values are are generally generally much much higher higher than than in in the the ICZ ICZ for for all all alloy alloy variants. variants. ThisThis can can be be due due to to the the significantly significantly coarser coarser grain grain size size and and the the absence absence of ofductile ductile polygonal polygonal ferrite ferrite in thein the microstructure. microstructure. The The absence absence of Ti of in Ti the in thealloy alloy leads leads to clearly to clearly higher higher DBTT DBTT in comparison in comparison to the to Ti-alloyedthe Ti-alloyed variant. variant. A significant A significant influence influence of moly of molybdenumbdenum for otherwise for otherwise identical identical alloys alloys cannot cannot be detected.be detected. In several In several of the of alloy the alloysystems, systems, the addition the addition of Nb ofresults Nb results in a slight in a decrease slight decrease of the DBTT, of the followedDBTT, followed by a marked by a marked increase increase for the highest for the highestNb content Nb contentof 0.1%. ofThe 0.1%. difference The di ffinerence DBTT in is DBTTmost obviousis most obviousbetween betweenalloys V4 alloys (MoNbTi) V4 (MoNbTi) and V5 (M andoNb). V5 (MoNb).Comparing Comparing the microstructures the microstructures of the CGZ of indicatesthe CGZ indicatesthat the Ti-free that the alloy Ti-free V5 alloyhas generally V5 has generally large prior large austenite prior austenite grain size grain (Figure size (Figure8), which8), decreaseswhich decreases somewhat somewhat towards towards the highest the highest Nb additi Nb addition.on. The The grain grain size size distribution distribution is is also also quite quite inhomogeneous.inhomogeneous. This This is congruentcongruent withwith thethe high high DBTT DBTT found found for for these these alloys. alloys. On On the the contrary, contrary, alloy alloy V4 V4has has much much smaller smaller and ratherand rather homogeneous homogeneous grain size.grain Yet, size. for Yet, alloy fo V4r containingalloy V4 containing 0.1%Nb individual 0.1%Nb individuallarge grains large are observed,grains are which observed, are probably which are the probably result of abnormalthe result grainof abnormal growth. grain The occurrencegrowth. The of occurrenceindividual largeof individual grains appears large grains to have appears a significant to have negative a significant influence negative on the influence DBTT. on the DBTT. The comparison of alloy V2-c with the corresponding industrial reference X65 steel reveals that the DBTT T50% is approximately 30 K lower in the HAZ sub-zones. In the ICZ the industrial steel has a DBTT of −130 °C as compared to −103 °C in alloy V2-c. In the CGZ the DBTT of the industrial steel is −65 °C whereas for alloy V2-c it is −32°C. If this difference, being related to steel cleanness was applicable to all other laboratory alloys, one can assume that the DBTT in the CGZ of industrially produced clean steels would be generally in the sub-zero temperature range. Appl. Sci. 2019, 9, 1847 10 of 17 Appl. Sci. 2019, 9, 1847 10 of 16

Figure 8.FigureEﬀect 8. of Effect Ti and of NbTi and micro-alloy Nb micro-alloy addition addition on themicrostructure on the microstructure in the CGZ in th (etchinge CGZ (etching with Le with Le Pera agent);Pera abnormalagent); abnormal prior austenite prior austenite grain size grain in alloy size in V4-d alloy (0.10%Nb) V4-d (0.10%Nb) indicated indicated by arrows. by arrows.

The3.3. comparison Post-Weld Heat of alloy Treatment V2-c with the corresponding industrial reference X65 steel reveals that the DBTT T50% is approximately 30 K lower in the HAZ sub-zones. In the ICZ the industrial steel During post-weld heat treatment, the strength of the material is lowered, thus, residual stresses has a DBTT of 130 ◦C as compared to 103 ◦C in alloy V2-c. In the CGZ the DBTT of the industrial introduced− to the weld zone can− relax by plastic deformation. The reduction of area during a hot- steel is 65 ◦C whereas for alloy V2-c it is 32◦C. If this diﬀerence, being related to steel cleanness tensile− test indicates how much deformation− the steel can withstand without cracking. Hot-tensile was applicable to all other laboratory alloys, one can assume that the DBTT in the CGZ of industrially testing was performed at 560 and 620 °C with steel that was prior subjected to the CGZ simulation produced clean steels would be generally in the sub-zero temperature range. heat cycle. Figure 9 shows the measured reduction of area as a function of the measured hot tensile strength. Taking the original tensile strength in the CGZ as the reference, heating to 560 °C reduces the tensile strength by 150–200 MPa. When heating to 620 °C the strength reduction is in the range of Appl. Sci. 2019, 9, 1847 11 of 16 Appl. Sci. 2019, 9, 1847 11 of 17 Taking the original tensile strength in the CGZ as the reference, heating to 560 °C reduces the tensile strength by 150–200 MPa. When heating to 620 °C the strength reduction is in the range of 230–3303.3. Post-Weld MPa. HeatAt both Treatment temperatures, the steels alloyed with molybdenum reveal a higher strength than Duringthose without post-weld it. heatIncreasing treatment, the the niobium strength co ofntent the materialsystematically is lowered, raises thus, the residual strength. stresses The reductionintroduced of to area the welddecreases zone canwith relax increasing by plastic tensile deformation. strength Thefor all reduction alloys at of both area temperatures. during a hot-tensile This decreasingtest indicates effect how is much stronger deformation at higher the steeltemper canatures. withstand At the without same cracking. hot-tensile Hot-tensile strength testinglevel, molybdenum alloyed steels allow a significantly higher reduction of area. The comparison of the was performed at 560 and 620 ◦C with steel that was prior subjected to the CGZ simulation heat cycle. reductionFigure9 shows of area the at measured 560 °C between reduction alloy of areaV2-c as (RA a function = 32%) and of the the measured industrial hot X65 tensile steel strength. (RA = 50%) reveals again the better performance of the latter.

(a) (b)

FigureFigure 9. ReductionReduction of area vs. hot-tensile strength ofof coarse-grainedcoarse-grained zonezone(CGZ) (CGZ) treated treated material material at at a

atest test temperature temperature of of (a ()a 560) 560◦ C°C and and (b ()b 620) 620◦ C.°C.

Finally,Taking thethe originaltoughness tensile behavior strength after in thethe CGZpost-w aseld the heat reference, treatment heating at 560 to 560 °C ◦forC reduces 30 min the is analyzedtensile strength in Figure by 150–20010a. One MPa.finds Whena strong heating correla totion 620 of◦C the the DBTT strength T50% reduction with the is measured in the range yield of strength.230–330 MPa.The majority At both of temperatures, the data can the be steelsfitted alloyedvery well with by molybdenuma linear relationship reveal aindicated higher strength by the dashedthan those line. without Only the it. Increasingdata of alloy the V1 niobium are deviating content systematicallyfrom this line. raisesThe yield the strength. strength Theof the reduction higher Nb-alloyedof area decreases variants with (0.05%Nb increasing and tensile 0.10%Nb) strength significantly for all alloys increases at both after temperatures. the heat treatment This decreasing by values ineff theect isrange stronger of 50 at to higher 175 MPa. temperatures. The Nb-free At alloys the same show hot-tensile little change strength in yield level, strength. molybdenum For the alloyed alloys containingsteels allow 0.02%Nb, a significantly the strength higher reductionincrease is of in area. the Therange comparison of 20 to 30 of MPa. the reductionIn the higher of area Nb-alloyed at 560 ◦C variantsbetween the alloy DBTT V2-c increases (RA = 32%)severely and by the 50 industrialto 100 K. In X65 the steelNb-free (RA variants,= 50%) the reveals DBTT again drops the by better 10 to 30performance K. In alloys of containing the latter. 0.02%Nb the DBTT increases in the range of 10 to 30 K. For comparing the data Finally,sets before the toughnessand after heat behavior treatment, after the Figure post-weld 7b is heatre-plotted treatment in Figure at 560 ◦9bC foron 30the min same isanalyzed scale as Figurein Figure 9a. 10It a.is Oneapparent finds that a strong Nb precipitation correlation ofmo theves DBTT the data T50% points with to the higher measured strength yield and strength. DBTT. ForThe the majority Nb-free of alloys, the data tempering can be fitted appears very to well lower by the a linear DBTT relationship while the yield indicated strength by theremains dashed nearly line. unaffected.Only the data Alloying of alloy of V1molybdenum are deviating does from not this seem line. to have The yielda particular strength influence of the higher on the Nb-alloyed toughness behaviorvariants (0.05%Nbafter this andheat 0.10%Nb)treatment. significantly Performing increases the same after post-weld the heat heat treatment treatment by valueson the inCGZ the of range the industrialof 50 to 175 X65 MPa. steel The results Nb-free in a alloysDBTT of show –55 little°C whereas change in in the yield comparable strength. Forlaboratory the alloys steel containing V2-c it is –0.02%Nb,21°C. the strength increase is in the range of 20 to 30 MPa. In the higher Nb-alloyed variants the DBTT increases severely by 50 to 100 K. In the Nb-free variants, the DBTT drops by 10 to 30 K. In alloys containing 0.02%Nb the DBTT increases in the range of 10 to 30 K. For comparing the data sets before and after heat treatment, Figure7b is re-plotted in Figure9b on the same scale as Figure9a. It is apparent that Nb precipitation moves the data points to higher strength and DBTT. For the Nb-free alloys, tempering appears to lower the DBTT while the yield strength remains nearly unaffected. Alloying of molybdenum does not seem to have a particular influence on the toughness behavior after Appl. Sci. 2019, 9, 1847 12 of 17 this heat treatment. Performing the same post-weld heat treatment on the CGZ of the industrial X65 Appl.steel Sci. results 2019, 9 in, 1847 a DBTT of –55 ◦C whereas in the comparable laboratory steel V2-c it is –21◦C. 12 of 16

(a) (b)

Figure 10. InfluenceInﬂuence of of post-weld post-weld heat heat treatment treatment (PWHT) (PWHT) on on strength strength and and toughness: toughness: ( (aa)) DBTT DBTT (T50%) (T50%)

vs. yield yield strength strength after a relaxation treatment at 560 ◦°CC for 30 min; ( b) DBTT (T50%) vs. yield yield strength strength before relaxation treatment (same data as Figure7 7b,b,same sametrendline trendline as as Figure Figure9 a).9a).

4. Discussion Discussion The present study has demonstrated that the infl influenceuence of various alloying elements on the HAZ properties requiresrequires a carefula careful analysis analysis of individual of individual metallurgical metallurgical effects such effects as grain such size, as hardenability grain size, hardenabilityand precipitation and status. precipitation In the investigated status. In the alloy invest variants,igated Nb alloy and variants, Ti are the Nb elements and Ti withare the potential elements for withprecipitation, potential Mo for andprecipitation, Mn reduce Mo the and carbon Mn activityreduce the and carbon increase activity the hardenability, and increase and the Sihardenability, increases the andcarbon Si increases activity in the austenite carbon [ 11activity,12]. Both, in austenite Mo and Nb[11,12]. are large Both, atoms Mo and having Nb are the tendencylarge atoms of segregatinghaving the tendencyto grain boundaries of segregating and exerting to grain solute boundaries drag [13 and,14 ].exer Titaniumting solute added drag up [13,14]. to the stoichiometric Titanium added ratio up with to therespect stoichiometric to nitrogen ratio contained with respect in the steel to nitrogen forms very contained temperature-stable in the steel forms TiN particles, very temperature-stable which exert grain TiNboundary particles, pinning which in hotexert austenite. grain boundary Niobium pinning carbide precipitatesin hot austenite. provide Niobium also grain carbide boundary precipitates pinning. provideHowever, also NbC grain precipitates boundary form pinning. at a lower However, temperature NbC precipitates than TiN and form typically at a lower niobium’s temperature precipitation than TiNis incomplete. and typically Thus, niobium’s in the interpretation precipitation of is the incomp presentlete. results, Thus, one in alwaysthe interpretation has to reflect of on the the present status results,of niobium, one always which in has thermodynamic to reflect on the equilibrium status of niobium, is described which by establishedin thermodynamic solubility equilibrium equations. is describedThe solubility by establishe of niobiumd solubility carbide, equations. where [Nb] and [C] are the alloy concentrations in mass percent, is describedThe solubility as a function of niobium of the absolutecarbide, temperature,where [Nb] and T, as [C] [15 ]:are the alloy concentrations in mass percent, is described as a function of the absolute temperature, T, as [15]: log [Nb][C] = 3.14 7753/T. (1) log [Nb][C] = 3.14− − 7753/T. (1) Accordingly, niobium additions of 0.02, 0.05 and 0.10% to the present 0.035% carbon steel dissolve Accordingly, niobium additions of 0.02, 0.05 and 0.10% to the present 0.035% carbon steel dissolve at at temperatures of 960, 1040 and 1120 C, respectively. The solubility of titanium nitride carbide is temperatures of 960, 1040 and 1120 °C,◦ respectively. The solubility of titanium nitride carbide is similarly described as [15]: similarly described as [15]: log [Ti][N] = 5.40 15790/T. (2) − log [Ti][N] = 5.40 − 15790/T. (2) In the present alloys, the nitrogen content [N] is 0.0070% so that according to Equation (2) TiN is stable Inbeyond the present the upper-temperature alloys, the nitrogen limit content of the austenite [N] is 0.0070% phase fieldso that for theseaccording alloys to being Equation around (2) 1450 TiN◦ isC. stableConsequently, beyond the at aupper-temperature peak temperature oflimit 900 of◦C the simulating austenite thephase inter-critical field for these sub-zone alloys both, being NbC around and 1450 °C. Consequently, at a peak temperature of 900 °C simulating the inter-critical sub-zone both, NbC and TiN particles are both stable. At a peak temperature of 1300 °C simulating the coarse- grained sub-zone, however, NbC particles should be dissolved whereas TiN particles are still stable. Therefore, TiN particles are available to control the austenite grain size in the CGZ simulation at 1300 °C. This is indeed obvious when comparing the CGZ microstructures of Ti-free alloys (coarse- grained) with those of Ti-added alloys (fine-grained). Niobium carbide particles will fully re-dissolve Appl. Sci. 2019, 9, 1847 13 of 17

TiN particles are both stable. At a peak temperature of 1300 ◦C simulating the coarse-grained sub-zone, however, NbC particles should be dissolved whereas TiN particles are still stable. Therefore, TiN particles are available to control the austenite grain size in the CGZ simulation at 1300 ◦C. This is indeed obvious when comparing the CGZ microstructures of Ti-free alloys (coarse-grained) with those of Ti-added alloys (fine-grained). Niobium carbide particles will fully re-dissolve during the up-heat cycle to 1300 ◦C. Solute Nb acts in the hot stage only by solute drag on austenite grain boundaries, which can contribute to finer grain size especially in combination with TiN particle pinning. This combined effect reflects an improved toughness. It is, however, also seen that for high Nb addition (0.10%) individual grains can grow abnormally large, leading to a significant decrease in toughness. In the as-rolled plates, a significant amount of niobium is expected to be in solution, out of equilibrium. This is due to the strong cooling after rolling and the low cool-stop temperature, preventing substantial precipitation of niobium as its precipitation kinetics is rather slow. Hence, any heat treatment of such an Nb super-saturated steel above temperatures of 500 ◦C and below the equilibrium dissolution temperature given by Equation (1) will effectively lead to Nb precipitation. The higher the temperature is within that range, the more complete will be the precipitation. Accordingly, Nb precipitation can be expected for the HAZ simulations at sub-critical and inter-critical peak temperatures. Differential analysis of the tensile data allows us to estimate the contribution by Nb precipitation strengthening. The result of this analysis is shown for the alloys with the highest Nb content for the two sub-critical temperatures in Figure 11a. The precipitation strengthening is generally larger at 800 ◦C than at 600 ◦C. Furthermore, the magnitude of the strengthening is with 80 to 100 MPa rather equal in all alloys. At 600 ◦C, the precipitation kinetics is slower and differences between the alloy systems become apparent. In alloys with an increased Mn content, the strength increase remains smaller and this effect is becoming even more pronounced when Mo is added. This can be due to the lower activity of carbon caused by Mn and Mo retarding the precipitation kinetics of Nb. On the other hand, when Si is reduced as in alloy V6, precipitation appears to be accelerated. It is assumed that the lower Si content allows for the better formation of nano-cementite particles [16], which serve as a precursor for the nucleation of NbC [17]. In the ICZ at a 900 ◦C peak temperature, niobium precipitation will occur even faster. Yet it is difficult to estimate the precipitation strengthening effect since simultaneously the entire matrix microstructure changes, also affecting the strength. The very pronounced grain refinement in the ICZ microstructure of Nb-alloyed variants certainly contributes to strength. This is caused for one part by the increased Ac1p temperature of the Nb alloyed steel vs. the Nb-free steel, which is a solute drag effect of Nb. The delayed partial transformation from ferrite to austenite enhances the nucleation rate for austenite grains. Secondly, the precipitation of Nb occurring simultaneous obstructs the growth of recrystallized ferrite and newly formed austenite grains. Re-distribution of carbon is more efficient due to shorter diffusion distances between grains in the refined ICZ microstructure. Therefore, the carbon content in newly formed austenite grains is more balanced and on average lower. This makes the formation of MA phase less likely, which is confirmed by the disappearance of the dual phase character (low YTR) that is clearly observed in the Nb-free steel. Molybdenum increases the hardenability of carbon enriched austenite so that a higher share of the hard phase is present in Mo-alloyed steels. This reflects in lower ICZ toughness values for the Mo-alloyed steels. However, increasing Nb content offsets the Mo effect via the carbon balancing explained before. The respective carbon-reduced austenite grains have a higher martensite-start temperature and will experience severe self-tempering during the down-cooling cycle. Accordingly, toughness improves rapidly in Mo-alloyed steels with increasing Nb content. The reduced Si content in alloy V6-d results in a particularly low Ac3p temperature (875 ◦C) for a reason that is not yet known. As a consequence, the peak temperature of 900 ◦C in the ICZ heating cycle causes an almost complete austenite transformation, preventing carbon enrichment in individual grains. Therefore, the formation of hard phases despite Mo alloying is unlikely. Indeed, the toughness of steel V6-d is on the same good level as found for the best Mo-free alloys. Appl. Sci. 2019 9 Appl. Sci. 2019,, 9,, 18471847 14 of14 16 of 17

120 240 220 100 200 180 MPa) MPa) ( ( 80 160 ng ng 140 heni heni 120

60 rengt rengt st st

on 100 on i i at at t t pi pi 40 80 Preci

Preci 60 20 40 20 0 0 V1 V2 V3 V4 V5 V6 V1 V2 V3 V4 V5 V6 V1 V2 V3 V4 V5 V6 V1 V2 V3 V4 V5 V6 V1 V2 V3 V4 V5 V6 Peak tem p. 600°CPeak tem p. 800°C Peak tem p. 1300°C PW HT 560°CSum

(a) (b)

FigureFigure11. 11. Precipitation strengthening strengthening effect eﬀect in in HAZ HAZ sub- sub-zoneszones for for alloy alloy variants variants containing containing 0.10%Nb: 0.10%Nb:

(a()a) Sub-critical Sub-critical zone zone at peakat peak temperatures temperatures of 600 of and600 800 and◦C; 800 (b )°C; coarse-grained (b) coarse-grained zone (peak zone temperature (peak 1300temperature◦C), precipitation 1300 °C), precipitation eﬀect of PWHT effect (560 of PWHT◦C/30 min.)(560 °C and / 30 sum min.) of and precipitation sum of precipitation strengthening (CGZstrengthening+ PWHT). (CGZ + PWHT).

WhenWhen heating heating up up to to the the CGZ CGZ peak peak temperature temperature of of 1300 1300◦C, °C, all all NbC NbC particles particles present present in in the the as-rolled as- platerolled will plate dissolve. will dissolve. The time The required time required for dissolving for dissolving such particlessuch particles having having typically typically a size a size below below 100 nm is100 estimated nm is estimated to be less thanto be oneless second.than one The second. TiN particles, The TiN however, particles, remainhowever, stable. remain Upon stable. down-cooling Upon niobium’sdown-cooling solubility niobium’s decreases solubility and decreases NbC should and re-precipitate.NbC should re-precipitate. Precipitation Precipitation of NbC does of not NbC easily occurdoes insidenot easily the austeniteoccur inside grain the in theaustenite absence grai ofn externally in the absence applied of plastic externally deformation. applied Theplastic lattice deformation. The lattice mismatch between the austenite matrix and the NbC phase, both having an mismatch between the austenite matrix and the NbC phase, both having an fcc structure, is around fcc structure, is around 25% [18]. Better nucleation conditions for precipitation are given in the 25% [18]. Better nucleation conditions for precipitation are given in the austenite grain boundaries. austenite grain boundaries. Furthermore, niobium, as well as carbon, segregates to the grain Furthermore, niobium, as well as carbon, segregates to the grain boundary resulting in a locally much boundary resulting in a locally much higher [Nb]∙[C] solubility product. Precipitation of Nb should higher [Nb] [C] solubility product. Precipitation of Nb should occur more efficiently during or after occur more· efficiently during or after the phase transformation from austenite to the bcc phase. the phase transformation from austenite to the bcc phase. However, due to slow kinetics, it is expected However, due to slow kinetics, it is expected to be incomplete. The precipitation strengthening effect tohas be been incomplete. estimated The as shown precipitation in Figure strengthening 11b by performi effectng a has differential been estimated analysis asof shownthe tensile in Figuredata of 11b bythe performing CGZ treated adi steels.fferential The analysismagnitude of theof precipit tensileation data ofstrengthening the CGZ treated ranges steels. between The 60 magnitude and 150 of precipitationMPa. The highest strengthening strengthening ranges is betweenfound in 60 alloy and 150V2-d MPa. containing The highest lower strengtheningMn and no Mo. is found The in alloyotherwise V2-d containingidentical Ti-free lower alloy Mn (V1-d) and no shows Mo. The clearly otherwise lower strengthening. identical Ti-free Likewise, alloy (V1-d) this is shows observed clearly lowerwhen strengthening.comparing alloy Likewise, V4-d and this isV5-d, observed where when also comparingthe Ti-free alloyalloy V4-dexhibits and V5-d,clearly where lower also thestrengthening. Ti-free alloy Possibly, exhibits the clearly larger lower austenite strengthening. grain size Possibly,in the Ti-free the largeralloys austenitecauses a more grain unequal size in the Ti-freedistribution alloys of causes Nb (due a more to boundary unequal segr distributionegation) and of Nb increased (due to diffusion boundary distances. segregation) and increased diffusionWhen distances. applying post-weld heat treatment to the CGZ treated steels, two main metallurgical effectsWhen can applyingbe expected, post-weld being a heat tempering treatment of the to thematrix CGZ and treated further steels, precipitation two main of metallurgical solute Nb. The eff ects cangrain be expected,size will not being be altered. a tempering Based of on the the matrix tensile and data, further the softening precipitation in the of matrix solute due Nb. to The recovery grain size willprocesses not be is altered. for all Basedalloy systems on the tensilebetween data, 40 and the 50 softening MPa. The in forthcoming the matrix due additional to recovery precipitation processes is forstrengthening all alloy systems is displayed between in 40Figure and 5011b. MPa. It is Thelargest forthcoming in the Ti-free additional steels and precipitation generally larger strengthening in the isalloys displayed with the in Figurehigher 11Mnb. addition It is largest and inthose the with Ti-free Mo steels alloying. and This generally means larger that more in the Nb alloys remained with the higherin solution Mn addition after the andCGZ those treatment with due Mo alloying.to the slower This precipitation means that kinetics more Nb in remainedthese particular in solution alloys. after Summing up the strengthening contributions after the CGZ and PWHT cycles results in rather similar the CGZ treatment due to the slower precipitation kinetics in these particular alloys. Summing up the values for all alloys with the exception of Si-reduced steel. strengthening contributions after the CGZ and PWHT cycles results in rather similar values for all The lower reduction of area observed in steels with increasing Nb content during the hot tensile alloys with the exception of Si-reduced steel. deformation can be due to in-situ precipitation. Freshly formed precipitates may obstruct dislocations generated during the hot deformation and reduce the ductility of the matrix. It is feasible that this effect is stronger in the grain interior while the area near to the grain boundaries might be depleted Appl. Sci. 2019, 9, 1847 15 of 17

The lower reduction of area observed in steels with increasing Nb content during the hot tensile deformation can be due to in-situ precipitation. Freshly formed precipitates may obstruct dislocations generated during the hot deformation and reduce the ductility of the matrix. It is feasible that this effect is stronger in the grain interior while the area near to the grain boundaries might be depleted from niobium as it has already precipitated after the CGZ cycle. This can lead to narrow areas in the vicinity of former austenite grain boundaries where plastic deformation will concentrate and cause rupture and a limited reduction of area. The observation that Mo alloying improves the hot ductility can be due to a retardation of in-situ precipitation and also to increased hot strength as Mo segregates to the grain boundary area. This should counteract the local concentration of plastic deformation. At the higher hot deformation temperature of 620 ◦C (Figure9b), the decrease of reduction of area with increasing Nb content is more pronounced since the precipitation kinetics is quicker. Yet also at this temperature, the precipitation retarding effect of Mo is evident. The understanding that in-situ precipitation of NbC during hot deformation is the main cause for reduced ductility allows designing practical solutions to alleviate problems when post-weld heat treating steel with higher Nb addition. Firstly, the post-weld heat treatment can be done at the lower side of the typical temperature range of 500 to 600 ◦C, preferably below 550 ◦C, slowing down the kinetics of in-situ precipitation. Secondly, molybdenum should be co-alloyed to niobium for additionally retarding the precipitation kinetics and stabilizing the niobium-depleted areas in the vicinity of grain boundaries. Titanium micro-alloying restricts the CGZ grain size, thus, providing improved toughness. Thirdly, welding should be executed using the multi-pass technique in a way that the CGZ of an earlier pass obtains sufficient tempering by a subsequent weld pass. In this way, precipitation is already being anticipated in a large part of the CGZ rendering less solute Nb for in-situ precipitation during PWHT. Since PWHT is typically applied to components fabricated from thick plate gages, multi-pass welding is usually applied. The parameters of the individual passes should be optimized to allow a maximum tempering effect on previous passes.

5. Conclusions Increasing addition of niobium raises yield strength and improves toughness simultaneously in the as-rolled state of the considered extra-low carbon steels. The combined addition of molybdenum and niobium provides the highest strength fulfilling the requirements for grade X80. Regarding the inter-critical heat-affected zone, niobium-free steels have the lowest yield strength and toughness. These steels exhibit a low yield-to-tensile ratio, which is caused by a fraction of martensite islands dispersed in recrystallized ferrite. Adding niobium in increasing amounts raises the yield strength as well as the yield-to-tensile ratio. The recrystallized ferrite grain size is being severely refined and martensite formation is being suppressed. Consequently, ICZ toughness is improving with increasing niobium content. Molybdenum addition increases the amount of hard phase in the inter-critical heat-affected zone leading to a decrease of ICZ toughness. The hard phase appears predominantly as tempered martensite. At higher niobium additions this toughness-reducing effect of molybdenum is being diminished. Lowering the silicon content further reduces the formation of hard phases. Regarding the coarse-grained heat affected zone, titanium-free steels show low CGZ toughness due to excessive grain coarsening. A small addition of titanium, typically in near-stoichiometric ratio to residual nitrogen present in the steel, counteracts pronounced grain coarsening. Increasing niobium content can further improve toughness in combination with titanium. At the highest niobium addition, however, individual grains showed abnormal coarsening in several alloys. The mechanism for this abnormal grain growth must be investigated in more detail. Increasing niobium content leads also to more pronounced precipitation strengthening reducing CGZ toughness. Thus, the actual toughness performance of the coarse-grained heat affected zone is controlled by the counteracting effects of grain refinement and precipitation. Appl. Sci. 2019, 9, 1847 16 of 17

Toughness in the coarse-grained heat aﬀected zone after post-weld heat treatment is controlled by the degree of in-situ precipitation of niobium that was prior dissolved during the CGZ heat cycle. The potential of in-situ precipitation increases naturally with the niobium content in the steel. However, by anticipating or delaying in-situ precipitation, PWHT embrittlement can be reduced to an acceptable level. This requires careful consideration of the integral alloy concept, the selected PWHT temperature and an appropriate multi-pass welding schedule. Molybdenum, particularly, was found to improve hot ductility, which helps to avoid stress relaxation cracking.

Author Contributions: H.M. evaluated and analyzed the data and wrote the paper. All interpretations of the observed effects reflect H.M.’s personal opinion. Funding: This work was funded in part by the International Molybdenum Association (IMOA), London, UK and by Companhia Brasileira de Metalurgia e Mineraçao (CBMM), Sao Paulo, Brazil. Acknowledgments: The author sincerely acknowledges the experimental work and data preparation by Franz Mayrhofer, Rupert Egger and Irina Eisheuer of Voestalpine Grobblech GmbH. Conflicts of Interest: The author declares no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

References

1. Gray, J.M. Application of Niobium-Molybdenum strengthening mechanisms in high strength linepipe steels. In Fundamentals and Applications of Mo and Nb Alloying in High Performance Steels; Mohrbacher, H., Ed.; TMS: Pittsburgh, PA, USA, 2014; Volume 1, pp. 121–134. 2. Mohrbacher, H.; Spöttl, M.; Paegle, J. Innovative manufacturing technology enabling light weighting with steel in commercial vehicles. Adv. Manuf. 2015, 3, 3–18. [CrossRef] 3. Hulka, K.; Gray, J.M. High Temperature Processing of Line Pipe Steels. In Niobium Science & Technology, Proceedings of the International Symposium Niobium, Orlando, FL, USA, 2–5 December 2001; TMS: Pittsburgh, PA, USA, USA, 2001; pp. 587–612. 4. Gray, J.M. Weldability of Niobium Containing HSLA Steel. WRC Bull. 1976, 213, 1–19. 5. Aronson, A.H. The Weldability of Columbium-bearing High Strength Low Alloy Steel. Weld. J. 1966, 45, S266. 6. Kirkwood, P. Niobium and Heat Affected Zone Mythology. In Proceedings of the International Seminar on Welding of High Strength Pipeline Steels, Araxá, Brazil, 27–30 November 2011; CBMM: Araxá, Brazil, 2011; pp. 23–64. 7. Bersch, B.; Kaup, K. Effect of Niobium on the Welding Behaviour of Thermomechanically Rolled Steels. 3R Int. 1983, 22, 484–492. 8. Wiebe, J.; Scheller, W.; Stallybrass, C. HAZ Properties of Nb Microalloyed High Strength Pipeline Steels. Steel Res. Int. 2017.[CrossRef] 9. De Meester, B. The Weldability of Modern Structural TMCP Steels. ISIJ Int. 1997, 37, 537–551. [CrossRef] 10. Morris, J.W. On the Ductile-Brittle Transition in Lath Martensitic Steel. ISIJ Int. 2011, 51, 1569–1575. [CrossRef] 11. Eckstein, H.J.; Fennert, M.; Ohser, J. Application of thermodynamic computations to the solution behaviour of niobium and vanadium carbonitrides. Steel Res. 1993, 64, 143–147. [CrossRef] 12. Koyama, S.; Ishii, T.; Narita, K. Effects of Mn, Si, Cr and Ni on the Solution and Precipitation of Niobium Carbide in Iron Austenite. J. Jpn. Inst. Met. 1971, 35, 1089–1094. [CrossRef] 13. Hoerner, M.; Speer, J.; Eberhart, M. Comparison of Ab-initio Solute-Boundary Binding Energies and Experimental Recrystallization Data in Austenite for Solute Nb and Other Elements. ISIJ Int. 2017, 57, 1847–1850. [CrossRef] 14. Raabe, D.; Herbig, M.; Sandlöbes, S.; Li, Y.; Tytko, D.; Kuzmina, M.; Ponge, D.; Choi, P.-P. Grain boundary segregation engineering in metallic alloys: A pathway to the design of interface. Curr. Opin. Solid State Mater. Sci. 2014, 18, 253–261. [CrossRef] 15. Kunze, J. Nitrogen and Carbon in Iron and Steel Thermodynamics; Wiley-VCH Verlag GmbH: Weinheim, Germany, 1990; pp. 178–206. 16. Kozeschnik, E.; Bhadeshia, H.K.D.H. Influence of silicon on cementite precipitation in steel. Mater. Sci. Technol. 2008, 24, 343–347. [CrossRef] Appl. Sci. 2019, 9, 1847 17 of 17

17. Hin, C.; Bréchet, Y.; Maugis, P.; Soisson, F. Kinetics of heterogeneous dislocation precipitation of NbC in alpha-iron. Acta Mater. 2008, 56, 5535–5543. [CrossRef] 18. Farjami, S.; Kubo, H. Elastic Energy Analysis of Carbide and Nitride-Type Precipitates in an Fe–Mn–Si–Cr Shape Memory Alloy. Mater. Trans. 2006, 47, 564–570. [CrossRef]

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