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Materials Characterization

journal homepage: www.elsevier.com/locate/matchar

Effect of niobium and phase transformation temperature on themicrostructure and texture of a novel 0.40% C thermomechanicallyprocessed steel

Vahid Javaheria,⁎, Nasseh Khodaieb, Antti Kaijalainena, David Portera

aMaterials and Production Engineering Unit, Faculty of Technology, University of Oulu, FinlandbDepartment of Chemical and Materials Engineering, University of Alberta, Canada

A R T I C L E I N F O

Keywords:TextureHot rollingMicrostructureMedium carbon steelNiobium

A B S T R A C T

Field emission scanning electron microscopy, electron backscatter diffraction (EBSD) and X-ray diffraction(XRD) have been employed to investigate the effect of niobium and phase transformation temperature on theevolution of microstructure and texture in a novel thermomechanically processed, medium‑carbon, low-alloysteel intended for slurry pipeline applications. Thermomechanical processing consisted of hot-rolling in theaustenitic region with deformation both above the recrystallization limit temperature and below the re-crystallization stop temperature. Immediately after rolling, specimens were directly quenched in water to twodifferent temperatures of 560 °C and 420 °C and subsequently furnace cooled from those temperatures to si-mulate the cooling of coiled strip on a hot strip mill. The microstructure of samples quenched to 560 °C mostlycomprised of upper bainite, whereas the samples quenched to 420 °C mainly consisted of lath-type lower bainite.The transformation texture of all samples at the mid-thickness position consisted of α, γ and ε-fibers with highintensities close to the transformed copper, transformed brass and rotated cube components. The addition of0.013 wt% Nb refined the microstructure and sharpened the texture. The texture of the small fraction of retainedaustenite present in the final microstructures indicated that the main bcc texture components result from thebrass and copper components in the parent austenite.

1. Introduction

For many years, the thermomechanical controlled processing(TMCP) of low-carbon steels has been widely employed in the pro-duction of large diameter pipes for the transportation of oil and gas inorder to improve the pipeline performance and lifetime. This isachieved by the enhancement of the yield strength and the toughness ofthe pipe material through grain refinement [1–3]. Grain refinement isachieved by controlling the austenite microstructure via the rollingparameters and controlling the subsequent austenite transformation viathe cooling path. In addition, microalloying with niobium is commonlyused to refine the grain size and optimize the processing. After slabreheating at a controlled temperature, usually in the range1100–1300 °C, TMCP involves controlled multi-stage hot deformation;rough rolling above the austenite recrystallization limit temperature(RLT) and then finish rolling below the recrystallization stop tem-perature (RST) [4–6] which is then followed by controlled cooling.During the rolling stage, repeated deformation and recrystallizationabove the RLT produces fine recrystallized austenite (γ) grains which

become flattened and elongated after deformation below the RST. Thislast rolling stage is sometimes referred to as pancaking. The elongatedand dislocated fcc austenite grains then transform to a fine bcc micro-structure (α) the details of which depend on the cooling pattern (eitherferrite, bainite or martensite). The bcc transformation products have acrystallographic texture which is related to the texture of the parentaustenite (γ). The nature and strength of this texture, in turn, dependson the amount of recrystallization above the RLT and deformationbelow the RLT [7].

When the finish rolling temperature is in the γ recrystallisationrange, a weak transformation texture is produced in α, but by in-troducing sufficient strain in the γ below the RST, a fairly sharp rollingtexture can be produced in γ, which translates into a sharp texture inthe α products [7]. In addition to microstructure and grain size, controlof texture is important with regard to the final mechanical properties asit affects the orientation dependence of the strength and toughness ofthe final product [8–11].

While many works have discussed the effects of TMCP on micro-structural and textural evolution in high-strength, low-carbon, low-

https://doi.org/10.1016/j.matchar.2018.05.056Received 27 March 2018; Received in revised form 28 May 2018; Accepted 29 May 2018

⁎ Corresponding author.E-mail addresses: Vahid.Javaheri@oulu.fi (V. Javaheri), Khodaie@ualberta.ca (N. Khodaie), antti.kaijalainen@oulu.fi (A. Kaijalainen), David.Porter@oulu.fi (D. Porter).

Materials Characterization 142 (2018) 295–308

Available online 30 May 20181044-5803/ © 2018 The Authors. Published by Elsevier Inc. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/BY-NC-ND/4.0/).

T

alloy steels like the linepipe steels API X80 and X100, e.g. [2, 8, 12–14],to our knowledge, there are only few studies on the effect of thermo-mechanical treatments on the microstructure and texture of mediumcarbon steel [15–17]. Therefore, the present study was aimed at in-vestigating the effects of thermomechanical processing and niobiummicroalloying on the microstructure and texture evolution in a novelmedium-carbon (0.40 wt%) low-alloy steel under consideration forslurry pipelines. We employed X-ray diffraction (XRD) and electronbackscatter diffraction (EBSD) analysis to explore the macro-texturaland micro-textural changes occurring during thermomechanical treat-ments of two steel compositions with 0 and 0.013 wt% Nb after twodifferent accelerated cooling patterns giving essentially isothermaltransformation of austenite to upper and lower bainite. FESEM andSTEM were utilized to characterize the microstructural components.

2. Experimental Material and Procedures

2.1. Material

The steel compositions used in this work were designed by the au-thors as a potential new slurry pipeline material [18]. To ensure lowimpurity levels, the steel was vacuum induction melted and cast into aclosed bottom iron mould with internal dimensions51×310×600mm. To reduce the prior austenite grain size andwiden the non-recrystallization temperature region during hot rolling,Nb was added as a microalloy element to one of the heats, see Table 1.

2.2. Hot Rolling Process

Two slab-shaped blocks (51×200×75mm) of each compositionwere cut from the main castings, reheated at 1200 °C for 3 h, and thenrolled on a laboratory rolling mill to a final thickness of 10mm. Rollingwas performed in two stages: 4 passes giving a total thickness reductionof 48% in the temperature range 1200–1100 °C followed by 4 passes

totalling a further 32% reduction in the temperature range 950–800 °C.The finish rolling temperature was chosen to be above the calculatedequilibrium A3 temperature, to ensure that the microstructure is fullyaustenitic at the start of the cooling process. Subsequently, the rolledstrips were directly quenched in water to 560 °C and 420 °C at ca.200 °C/s and then transferred to a furnace at the same temperature andallowed to furnace cool (ca. 0.01 °C/s) to simulate the coiling processon a hot strip mill. Temperature changes during the whole process wererecorded with thermocouples inserted into holes drilled into the mid-length, mid-thickness of the slabs. The thermocouple was also rolledtogether with the specimen enabling precise temperature control duringrolling and subsequent cooling. Fig. 1(a) shows a schematic diagram ofthe rolling and cooling schedules used. The codes used in the figure arelisted in Table 2. They indicate the different combinations of compo-sition (N niobium, no and Y niobium, yes) and simulated coiling tem-perature.

The aim of the two quench-stop temperatures (QST) was to achieveupper bainite and lower bainite during the subsequent, essentiallyisothermal, phase transformation. They were chosen on the basis of atemperature – time – transformation (TTT) diagram predicted usingJMatPro® software [19] for an austenite grain size of 25 μm (ASTM 8),see Fig. 1(b).

2.3. Dilatometry

Phase transformation temperatures and resultant microstructures foralloy 2 were determined using dilatometric measurements on a non-de-formed (unstrained) sample using a Gleeble 3800 thermomechanical si-mulator. For this purpose, a cylindrical sample was machined to a diameterof 6mm and a length of 35mm. The sample was heated in the Gleeble to950 °C at 5 °C/s, held for 30 s and then cooled to room temperature at5 °C/s. The change in diameter was measured during the whole process.

Table 1Chemical compositions (wt%).

C Si Mn Cr Ni Mo Nb Al N

Alloy 1 0.39 0.19 0.24 0.92 0.02 0.48 0.002 0.04 0.004Alloy 2 0.39 0.18 0.24 0.92 0.02 0.58 0.013 0.04 0.004

Fig. 1. a) Schematic presentation of thermomechanical rolling procedure. b) Coiling simulation process of Alloy 2 plotted on a TTT diagram calculated using JMatProsoftware.

Table 2Coding of the investigated materials.

Composition Coiling simulation temperature (°C)

420 560

Alloy 1 (without Nb) #N420 #N560Alloy 2 (containing Nb) #Y420 #Y560

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2.4. Texture and Microstructure

Microstructures were investigated using a light optical and a laserscanning confocal microscope (VK-X200, Keyence Ltd.) after metallo-graphic sample preparation and etching in 2% nital. In addition, higher

resolution microscopy was conducted using a 200 kV energy filteredscanning transmission electron microscope (JEOL JEM-2200FS EFTEM/STEM). From the mid-thickness of the samples, 3 mm diameter trans-mission electron microscopy samples were punched from ground0.1 mm thick wafers, which were then further ground to approximately0.08mm before electro-polishing at 25 V in an electrolyte consisting ofperchloric acid, butyl cellosolve, distillated water and ethanol.

A Zeiss Sigma field emission scanning electron microscope (FESEM)equipped with electron backscatter diffraction (EBSD) was employedfor microtexture characterisation. EBSD mappings were recorded usingan accelerating voltage of 15 kV, a working distance of 15mm, a tiltangle of 70o and a step size of 0.3 μm.

Macrotexture (XRD) measurements were performed on a RigakuSmartLab 9 kW X-ray diffractometer. All textures have been quantita-tively examined by measuring, from an area of 14× 24mm, three in-completes (110), (200) and (211) pole figures, which were used toconstruct orientation distribution functions (ODF) in three-dimensionalEuler space using LaboTex software. Samples were scanned over a 2θrange from 5° to 85° in the back-reflection mode with a step size of0.02° using radiation generated from a rotating Co Kα anode operatedat 40 kV and 135mA. All samples for the microtexture (EBSD) andmacrotexture (XRD) measurements were selected from the mid-thick-ness of the strips close to the mid-length position. The EBSD data werepost-processed with the MTEX toolbox [20, 21].

As it was difficult to satisfactorily reveal the parent austenite grainstructure using etching techniques, a computational reconstructiontechnique was applied to the EBSD results using Matlab supplementedwith the MTEX texture and crystallography analysis toolbox [22].Briefly, grain maps were initially assembled from the data sets with agrain boundary tolerance of 3–5°. Subsequently, the parent austeniteorientation map was reconstructed from this data with a two-step re-construction algorithm. Firstly, the orientation relationship betweenthe parent austenite and product ferritic phase was determined usingthe method proposed by Nyyssönen et al. [23] based on the Kurdju-mov–Sachs (KeS) relationship [24] (i.e., {111}γ//{110}α, ⟨110⟩γ//⟨111⟩α). In the second step, the grain map was divided into discreteclusters using the Markov clustering method [23] proposed by Gomesand Kestens [25]. The parent austenite orientation was then calculatedfor each cluster separately, resulting in a reconstructed austenite or-ientation map. The average misorientation between the reconstructedorientation for each cluster and the best fit for each individual grainwas approximately 2°, indicating a good fit for the reconstructed result.

Fig. 2. a) Dilatometry results for non-deformed material (Ø6×35mm) usingthe Gleeble 3800 and b) CCT diagram calculated using JMatPro software.

Fig. 3. Combined laser scanning and light optical images of the microstructures of the samples. Nital etching.

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The full details for the reconstruction procedure have been described byNyssönen [26].

3. Results and Discussion

3.1. Bainite Formation Temperature Range

Fig. 2 shows the dilatometry result and a continuous coolingtransformation (CCT) diagram for alloy 2 calculated using JMatProsoftware for an austenite grain size of 25 μm (ASTM 8), which wasappropriate for the reheated dilatometry specimen. The dilatometertrace indicates that for a cooling rate of 5 °C/s bainite starts to form at545 °C and martensite at 325 °C. The JMatPro software predicted thecorresponding temperatures to be respectively 551 °C and 339 °C for thecooling rate concerned. This indicates that the phase transformationalgorithms in the software are probably appropriate for the steel com-position concerned and gives confidence in the accuracy of the TTTdiagram of Fig. 1(b) used to choose the QSTs. Of course, the aboveresults apply to the continuous cooling of non-deformed material andthe bainite formation temperature range may be different when startingwith deformed austenite. However, the results suggest that the chosencooling paths, with QSTs of 560 °C and 420 °C, should result in fullybainitic microstructures.

3.2. Microstructure

3.2.1. Effect of Quench Stop Temperature (QST)Fig. 3 shows examples of the microstructures as seen in combined

laser scanning – light optical images. Both lath type bainite and gran-ular bainite is present for the QST of 560 °C, while a QST of 420 °Cresulted in mainly lath type bainite. Granular bainite appeared as small(< 10 μm) grains with lighter contrast than the lath type bainite.

In the samples containing Nb, the bainitic laths appeared to nucleateand grow from the prior austenite grain boundaries, whereas without

Nb, such a tendency was not apparent. Due to the complexity of bainiticmicrostructures, FESEM and EBSD analyses of lattice misorientationswas applied following the approach of Zajac et al. [27] to better identifythe microstructural constituents.

Image quality maps, inverse pole figures and misorientation angledistributions from 0° to 60° for the samples are shown in Fig. 4. It can beseen that the relative frequency of boundaries with low misorientationangles (< 15°) is much higher for QST 560 °C than QST 420 °C. Incontrast, the QST 420 °C microstructures (#N420 and #Y420) show ahigh incidence of high-angle misorientations (> 45°). By analysing themisorientation angles of different types of bainitic microstructures,Zajac et al. [27] concluded that upper bainite is characterized by amisorientation angle spectrum having a high proportion of low-angleboundaries (< 20°) and fewer high-angle boundaries (> 50°), exactlylike those of samples #N560 and #Y560 in Fig. 4(c), while the spectrumof lower bainite displays a maximum at around 50–60° like that inFig. 4(f). This indicates that the direct quenching to 560 °C and 420 °Cled to austenite transforming to the upper and lower bainite, respec-tively.

FESEM images of the microstructures of the samples #Y420 and#Y560 are presented in Figs. 5 and 6, respectively. Under isothermaltransformation conditions, bainite directly forms from austenite bynucleation and growth of bainitic ferrite [28, 29] accompanied by theformation of cementite either between the bainitic ferrite laths orwithin them. Figs. 5–7 show that different varieties of bainiticmorphologies were identified in the present samples as described in thefollowing paragraphs.

3.2.1.1. Plate Like Bainite (PLB). The majority of the microstructuralconstituents in both #Y420 and #Y560 samples, especially in QST420 °C samples, belonged to the plate-like bainite morphology, whichconsisted of irregular plate-like ferrite with fine cementite particlesinside the plates, see the in-lens image and schematic illustration inFig. 5(c). The STEM image in Fig. 7(b) shows that the cementite

Fig. 4. Image quality (a and d), inverse pole figure (IPF) map (b and e) and misorientation angle distributions (c and f). The colors in the IPFs show the crystal-lographic orientation of the rolling direction.

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particles inside the bainitic ferrite plates were faceted of variable size.Mandal et al. [30] recently reported similar observations for a 0.13% Cultra-high strength steel.

In secondary electron (SE) images, PLB could be seen to be lessetched than its surroundings and other microstructural components.The PLB in the lower bainite microstructures was finer than that inupper bainite most probably due to lower transformation temperatureinvolved.

3.2.1.2. Granular Bainite (GB). When, for any reason, the formation ofcementite is retarded, a carbide free granular bainite forms. The SEimage of Fig. 6(c) shows that the areas of granular bainite are deeplyetched and appear to be quite flat compared to the PLB. This is becausethe ferritic component of GB has a uniform composition andcrystallographic orientation. In the in-lens mode, it is very difficult todistinguish GB from the PLB, but PLB has a less equiaxed shape andcomprises an internal structure whereas GB has no internal

Fig. 5. a) FESEM secondary electron (SE) image of as-rolled microalloyed sample #Y420, b) In-lens image of highlighted red area of Fig. 5(a) following withschematic illustration of the lower lath like bainitic (LLLB) morphology and c) In-lens image of highlighted blue area of Fig. 5(a) following with schematic illustrationof the lower Plate like bainitic (LPLB) morphology. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of thisarticle.)

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substructure. The STEM image in Fig. 7(c) also clearly shows that thereare no internal features in the GB. This is the most importantcharacteristic of this kind of bainitic morphology.

3.2.1.3. Lower Lath Like Bainite (LLLB). Precipitation of cementitewithin the bainitic ferrite laths is characteristic of the lower lath-likebainite. According to Zajac et al. [27], in upper bainite, carbides mainlytend to precipitate on the boundaries between the laths, whereas in

lower bainite carbide precipitation occurs within the ferritic laths (asshown in Fig. 5(b)). A high-resolution STEM image of such amorphology is presented in Fig. 7(a).

3.2.1.4. Degenerated Upper Bainite (DUB). Continuous or intermittentlayers of retained austenite, distributed between the bainitic ferritelaths were found in sample #Y560. Such a morphology is known asdegenerated upper bainite (DUB) for both formed austenite on the lath

Fig. 6. a) FESEM microscopy secondary electron (SE) image of as-rolled microalloyed sample #Y560, b) In-lens image of highlighted blue area of Fig. 6(a) followingwith schematic illustration of degenerated upper bainitic (DUB) morphology and c) In-lens image of highlighted red area of Fig. 6(a) following with schematicillustration of the granular bainitic (GB) morphology. (For interpretation of the references to color in this figure legend, the reader is referred to the web version ofthis article.)

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boundary [27, 30] and interlath austenite [31, 32]. Two types ofretained austenite were observed within the bainitic ferrite laths:blocky austenite and austenite films. Similar observations have beenreported by Tirumalasetty et al. [33]. Fig. 7(c) shows a STEM image ofsmall islands of retained austenite, which show lighter contrast thancementite.

3.2.2. Effect of Nb on the MicrostructureFig. 8 shows the reconstructed prior austenite grain structure of

#N420 and #Y420 in the as-cast, slab reheated, and as-rolled condi-tions. For the slab reheated condition, as-cast material was reheated at1200 °C for 3 h and then immediately quenched in water to a fullymartensitic microstructure. To cover a larger area, a series of EBSDmeasurements were stitched together. The average grain size de-termined using the equivalent circle diameter method and the averagehardness value of each sample are also given in each figure. Theaverage hardness values are based on at least six indentations.

It can be seen that the sample containing Nb (#Y420) has a finerstructure and higher hardness than the sample without niobium(#N420) at all stages of the processing chain. These results are in linewith earlier reports [5, 34–41] that microalloying with Nb can refine

the microstructure of low-alloy steels through the whole productionprocess from solidification through slab-reheating to controlled-rollingand accelerated cooling resulting in a finer final microstructures andhigher strengths.

In ref. [18], the authors showed that, based on the simulation resultsfor alloy 2, the amount of Nb in the interdendritic region due to segre-gation during solidification could increase to 0.026wt%. According toequilibrium thermodynamic calculations (Fig. 9), Nb will be in solid so-lution at the soaking temperatures (1200 °C) in both dendritic and inter-dendritic region. On the basis of Fig. 9, it is unlikely that NbC will pre-cipitate during the rough rolling with perhaps the exception of theinterdendritic regions. However, Speer and Hansen [42] pointed out thatNb refines the austenite grain size by a solute drag effect. Maruyama andSmith [43] also showed that the onset of recrystallization can be preventedby small substitutional - interstitial atomic clusters. At sufficiently lowtemperatures, Nb(CN) precipitation can refine the austenite grain size byinhibiting grain boundary movement [36, 37] and thereby reducing thegrain growth between rolling passes [44]. Nb(CN) precipitation is alsoknown to inhibit subgrain boundary movement, thereby raising the RSTwhen compared to a Nb-free composition. This can lead to a greater de-gree of pancaking and deformation of the austenite without

Fig. 7. The scanning transmission electron image of identified bainitic morphologies. a) Lower lath like bainitic (LLLB) morphology, b) Lower plate like bainitic(LPLB) morphology, c) Degenerated upper bainitic (DUB) morphology and d) Granular bainitic (GB) morphology.

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recrystallization. This, in turn, leads to finer final α microstructures.The prior austenite grains seen on both RD-ND and TD-ND sections

of the hot-rolled materials exhibited a mixed distribution of coarse andfine austenite grains, which may be due some rolling reduction in thepartial recrystallization temperature regime. Nevertheless, on average,the Nb addition caused 23% grain refinement as seen on RD-ND sec-tions and 17% as seen on TD-ND sections which helped to produce 16and 9% increases in hardness, respectively [45]. The higher hardness ofthe samples containing Nb could be due to the presence of fine Nb(CN)precipitates or it could arise from a higher hardenability of the Nbadded samples arising from its 0.1 wt% higher Mo content (seeTable 1). It has to be noted that although Nb was added due to itsexpected benefit during thermomechanical rolling, the results showthat even during solidification Nb acts as a grain refiner too.

A finer austenite grain size is, of course, beneficial with respect to

the grain size of the final transformation microstructure. Bainitic ferritelath or plate thickness depends on three factors: 1) the strength of theaustenite at the transformation temperature, 2) the dislocation densityin the austenite and 3) the thermodynamic driving force for the γ→ αtransformation [46–48]. Finer bainite results from strong austenitecontaining a high density of dislocations and a large driving force. Bothaustenite strength and dislocation density refine the structure by in-creasing the resistance to the motion of the glissile interfaces, and thethermodynamic driving force refines the structure by increasing thenucleation rate. All three factors increase as the transformation tem-perature decreases, so a lower bainite transformation temperature leadsto a reduction in the thickness of the bainitic ferrite plates. Of the abovefactors, the presence of Nb is expected to increase the dislocationdensity of the austenite and thereby lead to some refinement of thebainite at given QST. This effect of Nb can arise in two ways, i.e. 1) an

Fig. 8. Reconstructed prior austenite grain structures. a and a′) As-cast material. b and b′) Reheated to 1200 °C for 3 h. c and c′) After hot rolling and quenching to420 °C as seen on RD-ND cross-sections. d and d′) After hot rolling and quenching to 420 °C as seen on TD-ND cross-sections. (Mean Equivalent Circle Diameter (ECD)grain sizes and hardness levels are given in each image. The green grain boundaries are the grain boundaries with different misorientation angle like thermal twinboundaries). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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increase in the RST giving the opportunity for more deformation and ahigher cumulated dislocation density in the no-recrystallization tem-perature regime, and 2) the retarding effect of Nb on the rate of dis-location recovery after the deformation before quenching [49].

The beneficial effect of Nb on the final mean effective ferritic grainsize is apparent in Fig. 10 for the samples #N420 and #Y420. Here theeffective grain size is taken as the ECD of grains surrounded by

boundaries with a misorientation angle of at least 15°. For the QST of420 °C, the addition of Nb reduced the ECD effective grain size (d90%)from 5.3 to 4.1 μm (d90% is the effective grain size at 90% in the cu-mulative grain size distribution). Comparing the grain size distributionsin Figs. 10(c) and (c′), it appears that the addition of Nb has an evengreater effect on the largest effective grain diameter recorded, which isaround 23 μm without Nb and 16 μm with Nb. This, almost one-third,reduction of the size of the coarse grains brought about by the Nb ad-dition may have greater consequences for the impact toughness of themicrostructure than the 22% reduction in the mean grain size: in thecase of direct quenched martensite, an inverse relationship between theCharpy V impact transition temperature and the size of the largestgrains in the effective grain size distribution has been found [50].

3.3. Texture

3.3.1. Macrotexture Evaluation (XRD Studies)In bcc metals, the most important texture components, in both the

deformed and annealed conditions, are found in the φ2= 45° section ofthe Euler space [51]. Fig. 11 shows such ODF maps as calculated fromthe X-ray diffraction pole figures as well as a schematic illustration ofthe ideal positions of the main components in the φ2= 45° section. TheRP value (error factor) for the calculation of the ODFs from the polefigures was 3.6% on average, which indicates good ODF reliability [52].The texture of the samples was typical of that found in hot bands offerritic steels when the transformation texture originates from a pan-caked austenite rolling texture, i.e. the main components are {113}⟨110⟩α,{111}⟨121⟩α, and {323}⟨131⟩α [53]. The characteristics of themain identified components and also their volume fractions are given inTable 3. The volume fraction was calculated as the fraction of grainswhose orientation lies within a±10° tolerance of each ideal orienta-tion.

Fig. 9. Phase diagram showing Nb solubility in austenite for the steel compo-sition 0.39C-0.20Si-0.25Mn-0.50Mo-0.90Cr-0.004 N (calculated using Thermo-Calc).

Fig. 10. a and a′) Inverse pole figure map showing the crystallographic orientation of the rolling direction. b and b′) High-angle grain boundaries in the bainiticmicrostructure and the equivalent circle diameter (ECD) mapping of each grain. c and c′) ECD grain size distribution of samples #N420 and #Y420, note the differentscales of the axes and color bar.

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The ODF sections show that the texture of the alloys is essentiallythe same for both QSTs, 420 °C and 560 °C, comprising mainly ND//⟨111⟩,i. e. the γ fiber, and RD//⟨110⟩, i.e. the α fiber. However, thevolume fractions of the main transformed components were greater insamples #N560 and #Y560 than in samples #N420 and #Y420.Moreover, the volume fractions of the main transformed components inthe samples containing Nb was a little greater than the sample withoutNb. In general, the sharpness of the texture and the relative intensitiesof the various components depend on the composition of the steel, the

amount of rolling reduction below the RST, the initial austenite grainsize, and the cooling rate during transformation [7]. Nonetheless, as allof these factors were similar for the both specimens, the sharper textureof the Nb microalloyed sample is most probably a result of a strongeraustenite texture due to the presence of Nb increasing the amount ofdeformation in the no-recrystallization temperature regime. It has beenwidely reported that the texture of Nb-microalloyed steels is con-siderably stronger than carbon steel as the precipitation of Nb(CN)suppresses the recrystallization of γ, thereby leading to a sharp γ rolling

Fig. 11. Centerline ODF sections at φ2= 45°. a) #N420, b) #Y420, c) #N560, d) #Y560 and e) the rolled bcc texture components in the φ2= 45° section.

Table 3The volume fraction of the main observed texture components and their Euler angles and orientations.

Symbol Component {hkl}⟨uvw⟩ Euler angles (°) Crystal orientation Component volume fraction

φ1 ∅ φ2 #N420 #Y420 #N560 #Y560

TC1 Transformed copper {112}⟨110⟩ 0 35 45 7.28 8.07 7.84 8.50

TC2 Transformed copper {113}⟨110⟩ 0 25 45 6.15 6.97 7.98 8.39

RC Rotated cube {001}⟨110⟩ 0/90 0 45 2.47 3.21 3.97 3.63

TB Transformed brass {111}⟨112⟩ 90 55 45 4.03 3.19 4.18 4.55

C Cube {001}⟨100⟩ 0 0 0 1.84 2.44 2.32 2.13

Gs Goss {110}⟨001⟩ 90 90 45 1.08 0.98 1.27 1.22

A1 – {525}⟨151⟩ 15 47 68 3.40 3.50 3.91 3.98

A2 – {323}⟨131⟩ 23 50 56 4.89 5.96 6.49 6.78

A3 – {110}⟨011⟩ 60 55 45 3.34 2.67 3.22 3.44

All other orientations 64.45 64.08 58.88 57.62

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texture [7, 16, 34, 54, 55].Fig. 12 presents macrotextural details of the α and γ-fibers, which

are the most important fibers for hot rolling under plain strain de-formation, and the ε-fiber. With regard to Fig. 12(a), in the α-fiber allthe samples have the maximum orientation intensity near {223}⟨110⟩αand {112}⟨110⟩α and the intensity of these orientations is almost thesame for all samples. Regarding Fig. 12(b), it is worth noting that the γ-fiber contained important recrystallisation components, especially closeto {111}⟨112⟩α and {111}⟨123⟩α. Fig. 12(c) shows that the orientationdensity of ε-fibers is mainly centered on {332}⟨113⟩α and {554}⟨225⟩αwhich are transformed components originating from the parent auste-nite brass component. Of course, due to symmetry, there is an intensitypeak at the beginning of the ε-fibers corresponding to that at the be-ginning of the α-fiber. It has been reported that, among the transfor-mation texture components found in controlled rolled low-carbon oiland gas pipeline steels, the {113}⟨110⟩α (TC2) component causes sig-nificant anisotropy in mechanical properties i.e. both strength andtoughness [3, 7, 8]. By contrast, less anisotropy in strength andtoughness is reached when the {332}⟨113⟩α component is strong. Thisindicates that the composition containing Nb may offer more isotropicmechanical properties than that without Nb as the intensity of the{332}⟨113⟩α component is increased by the addition of Nb. Comparing

the present steel with low-carbon steels like API X80 [8], API X100 [3],Nb HSLA steel [55, 56], and extra-low-carbon steel [57], despite of thedifference in the composition, rolling schedule, and cooling conditionwhich can mainly affect the final texture, the texture patterns aremostly the same but with the variation in the intensity. However, othertexture types have been also reported for low carbon steels like re-ference [58].

3.3.2. Microtexture Evaluation (EBSD Studies)The crystallographic characteristics of the samples were also ex-

amined using EBSD. The microtextural features of variant #Y420,which are representative of all the variants, are presented in Fig. 13 forthe RD-ND section. High-angle grain boundaries (HAGB) were definedas having a misorientation angle> 15° (black boundaries in Fig. 13(c)).Fig. 13 shows a combination of equiaxed and elongated grains with thetexture common to bcc rolled material. It can be seen that, apart fromthe rotated Goss (bottom left-hand corner), the analyzed area gives asimilar texture to that found using XRD and can therefore be consideredas representative of the specimen as a whole. However, a considerablefraction of the bainite has orientations outside the main componentslisted in Fig. 13(d). It can be seen that this is particularly true of thecoarser grained regions. Comparing Figs. 13(b) and (c) suggests that the

Fig. 12. Macro-textures of the samples. a) Alpha fiber, b) gamma fiber, c) epsilon fiber and d) schematic illustration of location α, γ and ε-fibers in the Euler Space.

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bainite that nucleated and grew intergranularly within the large pan-caked austenite displayed more random orientations while nuclei thatformed on the grain boundaries of finer grains are associated with themain components listed in Fig. 13(d). This would explain the somewhatsharper texture in the Nb microalloyed variants which had the finerprior austenite grain structures.

The final room temperature microstructures contained a smallfraction of retained austenite, and by assuming it to be representative ofall the austenite orientations present before quenching, it was possibleto determine the texture of the prior austenite grains. This is shown inFig. 14(a). It is in good agreement with the reconstructed texture ob-tained from EBSD data extracted from large areas of 100×500 μm forvariants #Y420 and #N420, as illustrated in Figs. 14(b–c).

By comparing Figs. 13 and 14, it could be concluded that the maintransformation texture components (transformed brass, transformedcopper, and rotated cube) originated from a pancaked austenite rollingtexture, with the characteristic components of copper ({112}⟨111⟩γ), Goss({110}⟨001⟩γ), and brass ({110}⟨112⟩γ), which is schematically illu-strated in the Fig. 15. This is to be expected since, when austenite is rolledbelow the RST, the deformation texture components are no longer re-moved by recrystallization. In the other words, the {332}⟨113⟩α compo-nent forms by transformation from the {110}⟨112⟩γ orientation present inthe deformed austenite, and the {113}⟨110⟩α component originates in asimilar way from the {112}⟨111⟩γ orientation present at the centerline ofthe rolled plate after the plain strain deformation.

4. Conclusions

The effect of Nb microalloying on the texture and microstructure ofa novel medium-carbon low-alloy wear resistant steel that was

thermomechanically rolled and direct-quenched to 560 °C and 420 °Cwith subsequent furnace cooling to room temperature has been studiedusing a combination of X-ray diffraction, FESEM/EBSD and STEM. Thefollowing conclusions can be drawn:

• The microstructure of samples with QST 560 °C mainly comprisedvarious upper bainite morphologies, i.e. conventional upper bainite,plate like bainite (PLB), granular bainite (GB) and degeneratedupper bainite (DUB), whereas the microstructure of QST 420 °Csamples were mostly lower plate-like bainite (LPLB) and lower lath-like bainite (LLLB).

• An addition of 0.013wt% Nb refines the prior austenite grain sizeand the effective grain size of the final bainitic microstructures, i.e.the size of grains bounded by high-angle boundaries. The addition ofNb was particularly effective in making the effective grain sizedistribution more uniform by reducing the size of the largest grainsby almost one third. The microstructural changes brought about bythe addition of Nb are reflected in an increase in the hardness of thefinal microstructures.

• Nb refined the prior austenite grain structure and sharpened itstexture thereby increasing the intensity of the bcc texture aftertransformation to either upper or lower bainite.

• The ODF of the alloys shows that the rolling reduction of the aus-tenite below the recrystallization stop temperature was high enoughto sharpen the desired {332}⟨113⟩α texture component, which isexpected to improve the isotropy of the mechanical properties.Moreover, macrotexture results show that the maximum intensitybelonged to the transformed copper component.

• Analysis of the retained austenite texture indicated that the mainbcc texture components are the result of transformation from the

Fig. 13. Variant #Y420. a) Inverse pole figure map of the bainitic structure showing the crystallographic orientations of the rolling direction. b) The reconstructedparent austenite grains. c) The location of the major components embedded in the high-angle grain boundary map. d) The color coding and volume fraction of eachcomponent. e) φ2=45o section of the ODF. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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brass and copper components in the parent austenite.

Acknowledgements

The authors are grateful for financial support from the EuropeanCommission under grant number 675715 – MIMESIS – H2020-MSCA-ITN-2015, which is a part of the Marie Sklodowska-Curie InnovativeTraining Networks European Industrial Doctorate programme. Theauthors would also like to thank EFD induction a.s, especially BjørnarGrande and John Inge Asperheim for their help and support during thisproject.

Data Availability

The raw/processed data required to reproduce these findings cannotbe shared at this time as the data also forms part of an ongoing study.

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