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Journal of Alloys and Compounds 689 (2016) 542e553

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Journal of Alloys and Compounds

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Precipitation strengthening in titanium microalloyed high-strengthsteel plates with new generation-thermomechanical controlledprocessing (NG-TMCP)

X.-L. Li a, C.-S. Lei a, X.-T. Deng a, *, Z.-D. Wang a, **, Y.-G. Yu a, G.-D. Wang a, R.D.K. Misra b

a State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang, 110819, Chinab Laboratory for Excellence in Advanced Steel Research, Department of Metallurgical, Materials and Biomedical Engineering, University of Texas at El Paso,El Paso, TX 79968-0521, USA

a r t i c l e i n f o

Article history:Received 24 June 2016Received in revised form19 July 2016Accepted 2 August 2016Available online 4 August 2016

Keywords:High strength low alloy (HSLA)Thermomechanical controlled processing(TMCP)NanoparticlesMicrostructureHigh resolution transmission electronmicroscopy

* Corresponding author.** Corresponding author.

E-mail addresses: [email protected] (X.-edu.cn (Z.-D. Wang).

http://dx.doi.org/10.1016/j.jallcom.2016.08.0100925-8388/© 2016 Elsevier B.V. All rights reserved.

a b s t r a c t

We elucidate here the strengthening mechanisms in titanium microalloyed low-carbon steels, whichwere rolled into plates of 12 mm thickness using a combination of thermomechanical controlled pro-cessing (TMCP) and ultrafast cooling (UFC). The ultrafast cooling combined with thermomechanicalcontrolled processing is referred by us as new generation (NG)-TMCP. Chemical phase analysis, small-angle X-ray scattering (SAXS) and high-resolution transmission electron microscopy (TEM) were usedto study the characteristics of nanoscale cementite precipitates and microalloyed precipitates. Besidesnanoscale TiC, cementite precipitates of size less than ~35 nm with high volume fraction were observedin Ti-microalloyed steel. Cementite with high volume fraction had a stronger precipitation strengtheningeffect than nanometer-sized TiC. The precipitation strengthening contribution of nanoscale precipitatesof different types and size was estimated together with solid solution strengthening and grain refine-ment strengthening contribution. The estimated yield strength was consistent with the experimentalvalue.

© 2016 Elsevier B.V. All rights reserved.

1. Introduction

Steel is an important structural material and strength is animportant property of any structural material. Strengtheningmechanisms have been studied for a number of years. Thermo-mechanical controlled process (TMCP) is a simple and cost effectivemethod to enhance the properties of ferritic steels [1], whichconsists of controlled hot rolling followed by controlled cooling.This involves complex interaction of chemical composition, tem-perature, deformation and different metallurgical phenomena. Thesuperior mechanical properties obtained via TMCP is a conse-quence of refinement of austenite, maximization of austeniteboundary area and density of deformation bands, such that thereare increased nucleation sites prior to transformation of austenite.The superior mechanical properties of low-carbon microalloyed

T. Deng), [email protected].

steels arise from a combination of refined ferrite grain size andprecipitation hardening.

In high-strength low-alloy (HSLA) steels containing micro-alloying elements such as V, Nb, and Ti, nanoscale V, Nb, and Ticarbides provide significant precipitation strengthening effect onthe addition of microalloying elements [2]. Steels industries arecurrently facing the challenging of reducing alloy costs in steelproducts, which negatively impacts the strength of steels. In orderto meet the requirements of reduction in alloy cost and maintainstrength, cementite in steels is viewed as a viable option to partiallyreplace microalloying elements because it is a common secondphase constituent in steels [3e5]. Fu and co-workers, observednanoscale cementite precipitates in Ti-microalloyed high strengthweathering steels during thin slab continuous casting and rolling,which were believed to play an important role in enhancing thestrength of steel [6]. Mao and co-worker studied Ti microalloyedweathering high-strength steel under laboratory and productionconditions [7] and found that TiC precipitation strengthening effectis significantly less than Fe3C precipitates. Other methods adoptedto obtain nanoscale cementites include cold deformation of pearl-itic steels after hot rolling [8] and static annealing after severe

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X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553 543

plastic deformation of a low carbon microalloyed steel [9]. How-ever, little work has been carried out to obtain nanoscale cement-ites in low-carbon steels plates by ultrafast cooling process aftercontrolled rolling. In order to clearly determine the possiblecontribution to strength, a detailed characterization of nanoscalecompounds is essential. Moreover, it is difficult to determine thevolume fraction of different types and size of nanoscale particles, totheoretically quantify their contribution to strength. Only qualita-tive information can be obtained by TEM.

The present study is focused on studying the effective contri-bution of nanoscale precipitates to strength during NG-TMCP ofmicroalloyed steel via chemical phase analysis, small-angle X-rayscattering (SAXS) and high-resolution transmission electron mi-croscopy (TEM). Additionally, the crystallography of different typesof precipitates is studied.

2. Materials and methods

2.1. Materials and thermo-mechanical processing

The chemical composition of the experimental steel is listed inTable 1. 0.08 wt% Ti was added based on the composition oftraditional Q345 grade. The steel was melted in a vacuum furnaceand cast into ingots of dimensions of 100 � 100 � 100 mm3, androlled into plates by TMCP process. Schematic illustration of TMCPschedule is presented in Fig. 1. The ingots were heated to 1250 �Cfor 2 h to completely dissolve titanium precipitates in austenite.After removal of the surface iron oxide, the ingots were rolled usingF450 mm rolling mill. The rolling process consisted of seven passeswith a total thickness reduction of 88%, and the final thickness was12 mm. Recrystallization of austenite commences at 1150 �C andconsisted of 3 passes of rough rolling. Austenite non-recrystallization region rolling started at 880 �C and final rollingtemperaturewas controlled at 860 �C. This process consisted of fourpasses. Next, the steels were ultra-fast cooled to isothermal holdingtemperaturewith different cooling rate and held for 20min. Finally,the steels were cooled in air to room temperature. The detailedprocessing parameters are described in Table 2, and the steels withdifferent processing parameters are marked as steel A and B inbelow.

2.2. Microstructural characterization

The steels were mechanically polished to mirror finish andetched with 4 vol% nital solution at room temperature. Micro-structural observations were carried out using a combination ofscanning electron microscopy (SEM, ZEISS ULTRA 55) and trans-mission electron microscopy (TEM, FEI Tecnai G2 F20). TEM speci-mens were produced by cutting slices from the steel plates,thinning them mechanically to 0.06 mm and electrochemically jetpolished in a solution of 9 vol% perchloric acid in ethanol at�30 �C.TEM experiments were carried out using a TEM equipped with anenergy-dispersive X-ray (EDX) spectrometer operating at 200 kV.Standard TEM techniques such as bright-field (BF) and dark-field(DF) imaging as well as selected area electron diffraction (SAED)were used to characterize the microstructure. The crystallographyand chemical composition of the precipitates were studied by highresolution transmission electron microscopy (HRTEM) and EDX.

Table 1Chemical composition of the tested steel used in this study (wt%).

C Mn Si Al Ti P S N O

0.15 0.98 0.28 0.02 0.08 0.015 0.005 27 ppm 48 ppm

Furthermore, the precipitates were electrochemically extractedfrom different steels. The phase constitution and size distributionwere studied using a PANalytical X'pert Pro MPD by using the smallangle X-ray scattering method, according to the test standard ISO/TS 13762.

2.3. Chemical phase analysis and SAXS analysis

It is quite difficult to determine the mass fraction or volumefraction of nanoscale particles in steels with an electron micro-scope, but the volume fraction of precipitates with different sizesare required for estimating the strength of steel. Chemical phaseanalysis and SAXS are effective methods to determine the chemicalcomposition, mass fraction, or volume fraction of nanoscale parti-cles with different sizes.

The procedure for chemical phase analysis was as follows:

(1) Electrolytic dissolution of steel sample to obtain electrolyzedpowders containing iron carbide, alloy carbide, sulfides andnitrides.

(2) Elimination of iron carbides, sulfides, and AlN to obtain alloycarbides and oxide.

(3) Removal of alloy carbides to get stable oxide.

After the above separation procedure, the electrolyzed powderswere dissolved and the content of Fe, Mn, N, and C was analyzed;and then the mass fraction of cementite based on formula(FeaMnb)3(CxNy) was calculated. According to GB/T 1322191 (ISO/TS13762 2001) the particle size distribution of precipitate powderswas analyzed by SAXS with a Kratky small-angle X-ray scattering(Rigaku Corporation, Tokyo, Japan). Using this approach, theanalyzed error was less than 10 pct.

2.4. Tensile and impact toughness tests

Standard cylindrical tensile test samples with a gage length of25 mm and diameter of 5 mmwere prepared from the heat treatedsteel plates transverse to the rolling direction. Tensile tests wereconducted at room-temperature to measure the yield strength,tensile strength and elongation using a SANS-5000 tensile testerwith a computerized tensile testing system at room temperature ata crosshead speed of 1 mm/min. Standard Charpy v-notch (CVN)

Fig. 1. Schematic diagram of NG-TMCP process at the laboratory scale.

Table 2Thermal mechanical processing parameters of the tested steels.

Steel Plate thickness Rolled inrecrystallizationtemperature region

Rolled in non-recrystallizationtemperature region

Cooling rate, �C/s Final temperature,�C Type of cooling

Start, �C Finish, �C Start, �C Finish, �C

A 12 1150 1096 889 864 64 580 Holding for 20, air coolingB 12 1150 1135 883 875 36 700 Holding for 20, air cooling

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553544

impact samples of dimensions 10 � 10 � 55 mm3 were prepared todetermine impact toughness at �20 and �40 �C using an Instron9250 impact tester.

3. Results

3.1. Microstructural evolution

Fig. 2 shows SEM and TEM micrographs corresponding to thesurface and mid-thickness in steel A. Fig. 2(a) and (c) illustrate themicrostructure of the surface and Fig. 2(b) and (d) describe themicrostructure of mid-thickness region of the plate. The micro-structure at both locations consisted of a fully bainite structurebecause of low finish cooling temperature and low cooling rate.TEM micrographs were indicative of upper bainite structure. Thewidth of bainite lath at the surface was less compared to the mid-thickness. Fig. 2(d) shows typical morphology of martensite-austenite (M/A) constituent of ~400 nm width. The M/A constitu-ents were systematically studied by Misra and co-wokers [10].During the slow cooling process, the undercooled austenite first

Fig. 2. (a, b) SEM and (c, d) TEM micrographs

transformed to ferrite and the untransformed austenite becomesrich in carbon. When C-rich austenite is cooled to a temperaturebelow Ms, it is partially or totally transformed to martensite andforms M/A constituent. In the present study, the M/A constituent insample A consisted of alternative layers of retained austenite andmartensite. The SAED pattern obtained from the area marked bywhite circle in Fig. 2(d) is illustrated in the inset in Fig. 2(d), sug-gesting Kurdjumov-Sachs (KeS) orientation relationship:½111�a==½101�g, (101)a//(111)g[11].

Fig. 3 is a schematic diagram of precipitation during differentstages in the experimental steel. During solidification and reheatingstage, TiN and Ti2CS are nucleated and redissolved, respectively.The typical morphology and EDS of TiN and Ti2CS are presented inFig. 4(a) and (b). It is well known that precipitation of TiN occurs intwo temperatures regimes: during solidification of steel and insupersaturated austenite after solidification. Coarse TiN pre-cipitates are formed at liquid iron temperatures of ~1540 �C duringsolidification [12]. In the d-g (austenite) transformation regime(around 1400 �C), the coarse TiN precipitates dissolve, leading to Tienrichment in d-g solid solution. At temperatures between 1400

of surface and mid-thickness of steel A.

Fig. 3. Schematic diagram of precipitation at different stages in the experimental steel.

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553 545

and 1200 �C, precipitation of TiN and Ti2CS precipitates takes placein supersaturated austenite [12]. Given that both TiN and Ti2CShave higher dissolution temperatures, they are expected to pin thegrain boundaries during hot rolling, leading to refinement of finalferritic grain size [13]. In the austenitization stage, both coarseningand redissolution of partial TiN and Ti2CS can occur. Duringdeformation of austenite, precipitation occurs at heterogeneousnucleation sites (e.g. dislocations) and/or at grain boundaries anddeformation bands. These precipitates are referred as strain-induced precipitates. In the ultrafast cooling stage, the growth ofstrain-induced precipitates is inhibited and precipitation in ferriteis promoted. Subsequently, during isothermal holding and aircooling stage, TiC and Fe3C are nucleated and coarsened in super-saturated ferrite or during g-a transformation, as shown in Fig. 4(c)and (d).

It is known that the precipitates can be clearly observed in thematrix, when the contrast of dislocations completely disappears.Thus, we tilted the specimen in TEM to satisfy the condition that

Fig. 4. SEM micrographs of (a) TiN and (b) Ti2CS formed during solidification and reheatingformed during the isothermal holding or air cooling stage.

the reciprocal vector of diffraction plane is perpendicular to theBurger's vector of dislocation [14]. Fig. 5(a) shows the morphologyand distribution of nanometer-sized carbides precipitated in lathbainite in the mid-thickness of the plate. Random dispersion ofcarbides of size range of ~2e10 nm was observed. RepresentativeHRTEM image of nanometer-sized carbide is presented in Fig. 5(b).As illustrated in the figure, the carbides were very fine and withinthe thickness of foil. This led to the development of Moir�e fringecontrast because overlapping carbide and ferrite lattices, whichprovide a clear contrast with the carbide and can be used for ac-curate measurement of the size of the carbide. The size of thecarbides in Fig. 5(b) were determined to be ~5.49 nm along thedirection of the Moir�e fringe (length) and ~5.26 nm perpendicularto the direction of the Moir�e fringe (thickness), and the aspect ratio(length/thickness) is close to 1. Thus, it can be concluded that themorphology of the carbidewas close to spherical. The average valueof length and thickness of carbidewas ~5.38 nm. In order to identifythe crystal structure and orientation relationship, the fast Fouriertransformed (FFT) diffractogram is presented in Fig. 5(c). Extradiffraction spots are related to iron oxide on the surface and doublediffraction from oxide and ferrite. The orientation relationship ofcarbide with respect to ferrite matrix is [110]carbide//[100]ferrite, andthe angle of the reciprocal vectors ð011Þferrite and ð101Þcarbide is~27.9�, which does not obey any specific orientation relationshipsuch as Baker-Nutting (BeN) and Nishiyama-Wassermann (NeW)orientation relationship. This result is inconsistent with previousstudies which indicated that carbides obey BeN or NeW orienta-tion relationship to have lowest lattice misfit and maintain co-herency at the interface with the ferrite matrix [15e22]. Along thezone axis½110�carbide==½100�ferrite, the double diffraction was foundto operate due to strong interaction betweenð220Þcarbide andð011Þferrite. The difference vector of the reciprocal vectorsfromð220Þcarbide and ð011Þferrite is illustrated in Fig. 5(c). Thespacing of a set of parallel Moir�e fringe was determined by thereciprocal of the magnitude of the corresponding difference vector,

stage, (c) Ti(C, N) formed during the austenite deformation or air cooling and (d) Fe3C

Fig. 5. (a) TEM, (b) HRTEM and (c) the corresponding FFT diffractogram of TiC precipitate; (d) BF image, (e) DF image and (f) SAD pattern of Fe3C at the mid-thickness of steel A.

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553546

and the Moir�e fringe is perpendicular to the difference vector. Thespacing for the produced Moir�e fringes is given by the followingequation [14]

D ¼d220carbide � d011ferrtie��

d220carbide � d011ferrtie�2 þ �d220carbide � d011ferrtie � ε2

��12 (1)

acarbide ¼ d�220

�carbide

� 2ffiffiffi2

p(2)

where D is the interplanar spacing of the Moir�e fringe, d220carbideand d011ferrite are the interplanar spacings of ð220Þcarbide andð011Þferrite,ε is the angle of reciprocal vectors ð220Þcarbide andð011Þferrite. The value of ε is 4.73� (0.0825 rad). The value ofd011ferrite was determined to be 0.202 nm by using the latticeparameter of ferrite (aferrite ¼ 0.287 nm). GGiven the measured D(0.643 nm) and ε (0.0825 rad), the lattice parameter of carbidedetermined from equation (2) was 0.433 nm.

It is known that Ti microalloying can result in precipitates in theferrite matrix. Surprisingly, not only TiC precipitates were observedbut also new type of fine precipitates was observed in Fig. 5(d). Tothe best of our knowledge, these precipitates have not been re-ported before in conventional HSLA steels plates. Thus, a detailedTEM study was carried out on these precipitates. The selected areadiffraction (SAD) pattern showed precipitates with [001] orienta-tion as seen in Fig. 5(f). It follows that these precipitates have anorthogonal structure. The lattice parameters a and b of the pre-cipitate were calculated to be 4.525 and 5.089. The precipitateswith these lattice constants and orthogonal structure has been

determined to be cementite with certainty. A DF image using the(210) reflection in the [001] projection of the precipitates is shownin Fig. 5(e). Here, all the precipitates that have same orientation arevisible. Furthermore, one can see in Fig. 5(b) that the size distri-bution of these cementites was in the size range of ~10e25 nm. Theshape of the observed cementites in Fig. 5(d) (identified by thecircle) and the streaking of precipitate reflections in the SAEDpattern in Fig. 5(f), suggested that these precipitates have disk-shaped morphology.

Fig. 6 shows the representative SEM and TEM micrograph cor-responding to surface and mid-thickness region in steel B. Fig. 6(a)and (c) indicated that the microstructure at the surface consisted of~3e8 mm fine polygonal ferrite (PF), and ~1e2 mm fine pearlite. Thepercentage of polygonal ferrite and pearlite were 89% and 11%,respectively. Fig. 6(b) indicated that the microstructure at mid-thickness of the plate consisted of ~5e10 mm coarse polygonalferrite, and ~1e3 mm fine pearlite. The percentage of polygonalferrite and pearlite were 85% and 15%, respectively. The micro-structure at the surface and mid-thickness mainly consisted ofpolygonal ferrite and pearlite. Compared to mid-thickness, thepolygonal ferrite was marginally refined because of higher under-cooling. Moreover, a higher fraction of pearlite was obtainedbecause of relatively low cooling rate, as shown in Fig. 6(a) and (b).The polygonal ferrite first forms at the original austenite grainboundary during phase transformation and rejects C to the sur-rounding austenite. The undercooled austenite was graduallyenriched with C, and pearlite formed during the slowing coolingprocess. Additionally, Fig. 6(d) illustrates fine lamellar cementitelocated between ferrite laths in pearlite with interlamellar spacingof ~308 nm.

Fig. 6. (a, b) SEM and (c, d) TEM micrographs of surface and mid-thickness of steel B.

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553 547

Fig. 7 shows precipitates in steel B. Two different morphologiesof precipitates were observed, interphase precipitation and randomprecipitates. The majority of interphase precipitates in rows were~5 nm in size and intersheet spacing was ~10e20 nm, as shown inFig. 7(a). The random precipitates in ferrite matrix were smallerthan ~10 nm and slightly larger than the interphase precipitates.The energy dispersive X-ray (EDX) spectrum of a representative12 nm precipitate in Fig. 7(b) is presented in Fig. 7(g), showing thatthe nanometer-sized carbide was TiC. According to the previousstudies, the interphase precipitation occurs at relatively highertemperatures, and that interphase precipitation can only occur insome ferrite grains and not in all [23e25]. In our study, interphaseprecipitation was observed only in some large ferrite grains whenthe specimen was ultrafast cooled to 700 �C. It is generally agreedthat the ledge velocity is inversely proportional to the isothermalaging temperature. The higher isothermal temperature reduces thedriving force for g/a transformation and ledge migration velocity,and increases the possibility of interphase precipitation [26].Representative HRTEM micrograph of interphase precipitation ofcarbides in ferrite is presented in Fig. 7(d). The crystal structure andthe orientation relationship were identified by the fast Fouriertransformed (FFT) diffractogram, as illustrated in Fig. 7(e). Extradiffraction spots are related to iron oxide on the surface and doublediffraction from oxide and ferrite. The interphase TiC precipitateshave NaCl-type crystal structure and obey the Nishiya-maeWassermann (NeW) orientation relationship with respect toferrite matrix: [110]carbide//[100]ferrite and ð011Þferrite==ð111Þcarbide[27]. By using inverse fast Fourier transformation (IFFT) for car-bides, the lattice images of the carbides along the zone axis

[110]carbide (Fig. 7(f)) were obtained, and the lattice parameters ofthe carbides was estimated to be 0.434 nm.

For fcc precipitates in a bcc matrix, the primary orientation re-lationships (OR) are BN, NW and Kurdjumov-Sachs (KS). Baker-Nutting (BN) relationship is {001}fcc//{001}bcc and fcc//bcc. On a given (001)fcc plane, there are two directions:[110] and ½110�. Because these two directions are perpendicular, asymmetry operation on fcc crystal with rotation of 90� about thepole of (001)fcc plane, would lead to [110] and ½110� being crystal-lographic equivalent. Thus, three variants of BN ORs can be derived.The NWORs can be obtained from a rotation of BN ORs. Each BN ORcan produce four variants or NW ORs, and therefore three variantsof BN ORs can bring about 12 variants of NW ORs. However, in thepresent study for TiC carbide interphase precipitation, only a singlevariant of NW OR with respect to ferrite matrix was observed. Theresult is consistent with interphase precipitation of (Nb, Ti)C [27].

In recent years, Zhao et al. and Fu [6,28] studied the distributionof nanoscale cementite in 0.15Ce1.5Mn ferrite-pearlite steel, andobserved less than 100-nm-diameter spherical cementite particles.The distribution of these particles was not homogeneous and thecontent was high in some regions. In order to study whether theaddition of titanium would affect the formation of cementite, TEMstudies were conducted. The results indicated that cementite andTiC precipitated simultaneously in the same ferrite grain, as shownin Fig. 7(c). The SAED pattern (Fig. 7(c)) confirmed that precipitatesof size range ~10e40 were cementite, with the zone axis of [111].The representative EDS spectrum of these precipitates in Fig. 7(c) ispresented in Fig. 7(h), showing that the relatively larger pre-cipitates are iron carbides.

Fig. 7. TEM micrograph of (a) interphase TiC, (b) random TiC, (c) Fe3C, (d) HRTEM image, (e) corresponding FFT diffractogram, and (f) IFFT lattice image interphase TiC; Repre-sentative EDX spectra of precipitates (g) in Fig. 7(a) and (h) in Fig. 7(c).

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553548

3.2. Mechanical properties

The mechanical properties including yield strength, tensilestrength, percentage elongation, and yield ratio obtained from

Table 3Mechanical properties of the tested steel.

Steel Yield strength, MPa Tensile strength, MPa

A 650 750B 590 718

tensile tests of Ti-microalloyed steel are listed in Table 3. The data inTable 3, is the average value of at least three specimens. In steel A,the yield strength, tensile strength, elongation and yield ratio were650MPa, 750MPa,19.2%, and 0.87, respectively; and in steel B were

Elongation, % Yield ratio Impact toughness (V-notch,J)

�20 �C �40 �C19.2 0.87 93.4 65.718.3 0.83 24.6 17.3

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553 549

590 MPa, 718 MPa, 18.3%, and 0.83, respectively.The CVN impact energy at subzero temperature (�20

and �40 �C) for steel A and B are shown in Table 3. For steel A, theCVN impact energy at �20 and �40 �C was 93.4 and 65.7 J,respectively, while for steel B, 24.6 and 17.3 J. The results indicatedthat the low-temperature impact toughness strongly depends onthe cooling rate and the finish cooling temperature. The fracturesurface of CVN impact samples tested at �40 �C were studied andtheir secondary electron fractograph are shown in Fig. 8. Fig. 8(a)shows fracture morphology at low magnification for steel A.Fig. 8(b) and (c) are high magnification images taken from the re-gion marked by marked by blue and pink rectangle in Fig. 8(a).Fig. 8(b) shows numerous small and large deep dimples, indicativeof good low-temperature toughness, while in Fig. 8(c) themorphology of fracture surface mainly consisted of quasi-cleavagefacets. In addition, some river-like marking on quasi-cleavage facetwith some dimples were observed. Fig. 8(e) shows the fracturemorphology of steel B. Fracture surface was predominantly cleav-age fracture with a small quantity of cracks. Based on SEM analysisof both steels, as shown in Fig. 8(d) and (f), large particles werepresent. Thus, high stress concentration can occur at the interface

Fig. 8. The secondary electron fractographs of CVN impact sa

due to difference in hardness between ferrite matrix and cementite[29,30], and when the stress is high, microvoids can nucleate. It isknown that these microvoids usually nucleate at the second phaseparticles, i.e. fineM/A constituent, or carbides, by de-bonding of theparticle-matrix interface or by particle fracture [31]. In addition, itwas found that the size of particles at the nucleation sites plays asignificant role in microvoid initiation [32]. Based on Griffith theory[33], for coarse particles in steel B, the crack initiation energy mustbe low, which is the reason for low toughness of steel B comparedto steel A. The difference in mechanical property is due to differentcooling schedules. It can be seen that a good combination ofstrength and ductility can be obtained with cooling rate and finishcooling temperature of 64 �C/s and 580 �C, respectively.

4. Discussion

4.1. Chemical phase analysis and SAXS

Chemical phase analysis and SAXS analysis were carried out onsteel A and B. Table 4 summarizes the structural parameters ofprecipitates obtained by X-ray diffraction for the extracted particles

mples tested at �40�Cof (aed) steel A and (e, f) steel B.

Table 4Structural parameters of precipitate.

Phase structure Lattice constant, nm Crystal system

M3(C,N) a0 ¼ 0.4523e0.4530, b0 ¼ 0.5088e0.5080, c0 ¼ 0.6743e0.6772 OrthorhombicTi2CS a0 ¼ 0.3210e0.3240, c0 ¼ 01.1203e1.1308, c/a ¼ 3.49 HexagonalTiC a0 ¼ 0.431e0.433 Face centered cubicTi(C,N) a0 ¼ 0.425e0.427 Face centered cubic

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553550

by electrolysis. From the X-ray results, we can say that the pre-cipitates were mainly composed of M3(C,N), Ti2CS, TiC and Ti(C,N).But it was difficult to distinguish between TiC and Ti(C, N), becauseof similarity in lattice constants.

Table 5 shows the mass fraction of precipitates in steel. FromTable 5, it can be seen that the highmass fraction of the precipitatedphase was M3C cementite-based compounds, followed by Ti(C, N).The size distribution of MC type and M3C type of steel A and B arepresented in Fig. 9. It may be noted that the MC carbide massfraction changed to a small extent in steel A and B, while the M3Cprecipitates in steel A and B were significantly different in the sizerange of ~1e36 nm. The mass fraction of M(C, N) particles of sizeless than 36 nmwas significantly less than M3(CxNy) of similar sizein steel. In steel A, the volume fraction of cementite with size lessthan ~36 nm was ~49 times of TiC of similar size and the corre-sponding yield strength increment of cementite was ~3.3 times ofTiC. In steel B, the volume fraction of cementite with size less than~36 nmwas ~4.8 times of TiC of similar size and the corresponding

Table 5Mass fraction of elements in M3C and MC in steel, Mass percent.

Steel Phase Mass fraction in steel of e

A (Fe0.984Mn0.016)3C Fe1.4806

Ti(C0.654N0.346) Ti0.0254

Ti2CS Ti0.0173

AlN Al0.0011

B (Fe0.974Mn0.026)3C Fe1.5220

Ti(C0.768N0.232) Ti0.0457

Ti2CS Ti0.0169

AlN Al0.0005

Note: C* represents calculated value.

Fig. 9. Size distribution of (a) MC-type precipitate a

yield strength increment of cementite was ~1.7 times of TiC.

4.2. Role of precipitates in strengthening of the steels

An overall objective of TMCP is to enhance the strength of steelby increasing the strength of the ferrite phase through precipitationhardening. It is therefore important to understand the contributionof the above mentioned precipitates toward effective strength-ening. Both TiN and Ti2CS can be very effective in pinning austenitegrain boundaries [12,13,34] at elevated temperatures during thehot rolling process. This results in reducing the ferrite grain size(grain refinement strengthening) and leading to enhancement ofyield strength and toughness. Precipitation strengthening is astrengthening method in which the interaction between dispersedfine precipitation phases and the dislocations in steel obstruct themovement of dislocations, thus increasing the strength of the steel.There are two well-known mechanisms by which the precipitatescan retard the motion of dislocation. One is called the shearing

lements Mass fraction of phase in steel

Mn C* 1.61280.0243 0.1079C* N 0.03220.0042 0.0026S C* 0.02530.0058 0.0022N 0.00170.0006Mn C* 1.67400.0400 0.1120C* N 0.05760.0088 0.0031S C* 0.02530.0021 0.0057N 0.00080.0003

nd (b) M3C-type precipitate in steels A and B.

Table 7The solid solution elements content in the experimental steels (wt.%).

Steel C N Mn Al Ti Si P

A 0.0357 0.0018 0.95 0.019 0.0373 0.28 0.015B 0.0271 0.0016 0.93 0.02 0.0174 0.28 0.015

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553 551

(cutting) mechanismwhere dislocations cut the particles; the otheris called the bypass (looping) or Orowan mechanism where thedislocations do not cut the particles, but bypass them, forming adislocation loop around the particles. The mechanism that requiresleast energy will be the relevant one for obstructing the dislocationmotion.

Considering that iron carbides are generally hard and resistcutting, the Orowan mechanism is most likely the dominantmechanism. The precipitation strengthening effect of fine pre-cipitates in steel can be calculated using the simplified equation (3)proposed by Gladman et al. [1] based on Ashby-Orowan's revisedmodel. When the average diameter of particles is greater than40 nm, the contribution of precipitation strengthening is notsignificant.

s ¼ 10:8ffiffiffif

pd

lnð1630dÞ (3)

In equation (3), s represents the precipitation strengtheningincrement in MPa, f is the volume fraction of the carbides and d isthe diameter of the carbide in mm.

When the diameter of precipitates is very small, and the inter-face tension between them and the matrix is small, the precipitateis coherent or semicoherent [35]:

s ¼ MtP ¼2� 1:1ffiffiffiffiffiffiffiffiffi

2AGp � g

3 =2

b2� d1

=

2f1 =2 (4)

where, tP represents the shear stress caused by the dislocationscutting the particles in MPa; A ¼ 1/2pKln(d/2b), represent thedislocation line tension function; K ¼ (1 � n) for edge type dislo-cation, K ¼ 1 for screw dislocation, and 1/K ¼ ½(1 þ 1/1�n) formixed type dislocation; n is the Poison's ratio and equal to 0.291; bis the absolute value of the dislocation Burger's vector and is equalto 0.248 nm; G is the shear elasticity modulus and equal to80,650 MPa; g is the interface energy between precipitates andmatrix and equal to 0.5e1 J/m2; d represents the second-phaseparticle diameter in mm; f represents the volume fraction of theprecipitates; and M represents the average Schmid orientationfactor and equal to 2 for bcc iron.

From equations (3) and (4), it can be seen that when the second-phase precipitate strengthening is shearing (cutting) type, itsstrengthening effect is proportional to the half power of particlesize d. When it is a bypass mechanism, its strengthening effect isproportional to half power of the volume fraction and roughlyinversely proportional to particle size d; that is, the strengtheningeffect of the shearing (cutting) mechanism increases withincreasing particle size, while the strengthening effect of the bypassmechanism decreases with increase of particle size. The criticaltransformation size dc can be calculated by the numerical solution

Table 6Contribution of Fe3C and TiC to the yield strength of steel A and B.

Steel Diameter range, nm TiC

Volume fraction, Pct Yield strength increment, MPa

A 1e5 0.0198 47.45e10 0.0013 12.810e18 0.0016 10.218e36 0.0028 7.4P

0.0255 77.8B 1e5 0.0284 57.1

5e10 0.0064 27.110e18 0.0015 9.718e36 0.0036 8.4P

0.0399 102.3

of equation (5): [5,36].

dc ¼ 0:209Gb2

Kgln�dc2b

�(5)

The meaning of symbols in equation (5) is same as mentionedabove. The critical transformation size of the precipitate dependson the properties of dispersed particles; the smaller the size,smaller is the interface energy between the precipitates and thematrix and larger the dc. Different types of precipitates havedifferent dc value.

For HSLA steels, it is generally believed that the precipitationstrengthening phase is nanoscale carbides of microalloying ele-ments. In our study, the precipitation strengthening is derived fromnanoscale iron carbide and TiC. In considering the contribution ofprecipitates to yield strength, the combined contribution of pre-cipitates with different types and sizes based on the bypassmechanism and shearing mechanism should be taken into account:

ssp ¼Xni¼1

ssp1i þXni¼1

ssp2i

¼Xni¼1

10:8

ffiffiffif

pd1i

lnð1630d1iÞ þXni¼1

2� 1:1ffiffiffiffiffiffiffiffiffi2AG

p � g3 =2

b2

� d1=

22i f

1 =2; d1i� dc � d2i (6)

where i represents the nanoscale precipitate; ssp1i represents theprecipitation strengthening contribution to the yield strength ofsteel based on bypass mechanism; and ssp1i represents the pre-cipitation strengthening contribution to the yield strength of steelbased on shearing mechanism.

The dc of TiC and Fe3C were estimated to be ~1.5e6 nm and~4.7e10 nm. The precipitation strengthening effect of TiC pre-cipitates with different sizes was calculated based on the bypassmechanism and for nanoscale iron carbide with less than 10 nmsize, the precipitation strengthening effect of precipitates wascalculated based on the shearing mechanism, and for nanoscaleiron carbides with size larger than 10 nm, the precipitationstrengthening effect was calculated based on the bypass mecha-nism. The calculated results are listed in Table 6. It can be seen that

Fe3C Total increment, mpa

Volume fraction, Pct Yield strength increment, MPa

0.712 144.6 331.10.074 150.252 51.20.209 42.51.247 253.30.123 118.7 279.40.0014 12.40.0014 9.50.0654 36.50.191 177.1

Table 8Comparison of calculated yield strength values with those actual measured.

Steel Grain size d, mm Calculated yield strength Values, MPa Actual measured ss, MPa

Grain refinement strengthening Solid solution strengthening Precipitation strengthening ss

A 5e6.5 235.3e268.3 78.6 331.1 645e678 650B 5.5e7.1 223.9e255.8 76.2 279.4 579.5e611.4 590

X.-L. Li et al. / Journal of Alloys and Compounds 689 (2016) 542e553552

the corresponding yield strength increment of cementite is 3.3times of TiC in steel A, while it is 1.7 times of TiC in steel B.

4.2.1. Comprehensive strengthening mechanismYield strength is the minimum stress at which dislocations

move and plastic deformation occurs. The greater the resistance todislocation motion, the greater is the yield strength. It is generallyagreed that the strengthening mechanism consists of solid solutionstrengthening, grain refinement strengthening, precipitationstrengthening and dislocation strengthening. While, consideringthat the original density of dislocation is related to the pinning ofthe precipitated secondary phase particles; the dislocationstrengthening can be ignored. According to Takaki [37], ultrafinegrain strengthening and dislocation strengthening or precipitationstrengthening cannot be added together. The underlying reason isthat when the value of ultrafine grain strengthening is calculated,the width of bainite or martensite lath is adopted as the grain size.Accompanying bainite or martensite transformation, a largeamount of secondary phase particles and dislocations pinned bynanoscale precipitates in the steel, whichmeans that ultrafine grainstrengthening includes or involves dislocation strengthening andprecipitation strengthening. If dislocation strengthening or pre-cipitation strengthening is added with ultrafine grain strength-ening, it will be added twice. Thus, for low-carbon steel, the yieldstrength of steel equals the sum of solid strengthening, grainrefinement strengthening, and precipitation strengthening, and isgiven by

sy ¼ sg þ ss þ ssp¼ 600D�1=2 þ f46½C� þ 37½Mn� þ 83½Si� þ 59½Al� þ 2918½N�

þ 680½P� þ 80:5½Ti�g þ ssp(7)

where, sy, sg, ss, ssp and represents the yield strength, the increasedstrength due to solid solution strengthening and the increasedstrength due to solid solution strengthening in MPa, respectively; Dis the average grain size in mm; [X] is the weight percent of alloyingelements, the elements in ferrite can be determined to be originalchemical composition removing carbide forming elements, asshown in Table 7. The components of the yield strength and thetotal yield strength calculated by equation (6) are presented inTable 8. The value calculated by equation (7) agrees with themeasured values. The contribution of precipitates to yield strengthcan be greater than ~300 MPa, which clearly indicates that pre-cipitation strengthening primarily contributed to the total yieldstrength.

5. Conclusions

(1) The microstructure of steel A with the finish cooling tem-perature of 580 �C and cooling rate of 64 �C/s mainly con-sisted of bainite at surface and mid-thickness because of thelow finish cooling temperature and the low cooling rate, andthe M/A constituent obeyed the Kurdjumov-Sachs (KeS)

orientation relationship. In contrast, in steel B, only ferriteand pearlite were observed.

(2) Random dispersion of carbides was observed in the sizerange of ~2e10 nm in steel A, and the random precipitatesdid not obey Baker-Nutting (BeN) and Nishiyama-Wassermann (NeW) orientation relationship. While insteel B, both random and interphase precipitates wereobserved. The interphase TiC precipitates had NaCl-typecrystal structure and obeyed the NeW orientation relation-ship with respect to ferrite matrix.

(3) Besides nanoscale TiC, cementite precipitates of size lessthan ~35 nmwere also observed in Ti-microalloyed steel. Themeasurement of volume fraction and size distribution ofnanoscale cementite and microalloyed carbide in steel werecarried out using chemical phase analysis and SAXS. Thevolume fraction of Fe3Cwas significant higher than Ti(C, N) ofsimilar size range.

(4) A good combination of strength and ductility can be obtainedat cooling rate and finish cooling temperature of 64 �C/s and580 �C. The nanosacle cementite of size less than ~35 nm hadremarkable precipitation strengthening effect, and thestrength increment was 253 MPa, which were ~3.3 times ofTiC. The calculated yield strength was consistent with theexperimental value.

Acknowledgements

The research was supported financially by National ScienceFoundation of China (Grant No. 52134002, 51504064, 51474064)and National Key Research and Development Program2016YFB0300601. RDKM gratefully acknowledges support fromUniversity of Texas at El Paso, USA.

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Precipitation strengthening in titanium microalloyed high-strength steel plates with new generation-thermomechanical contro ...1. Introduction2. Materials and methods2.1. Materials and thermo-mechanical processing2.2. Microstructural characterization2.3. Chemical phase analysis and SAXS analysis2.4. Tensile and impact toughness tests

3. Results3.1. Microstructural evolution3.2. Mechanical properties

4. Discussion4.1. Chemical phase analysis and SAXS4.2. Role of precipitates in strengthening of the steels4.2.1. Comprehensive strengthening mechanism

5. ConclusionsAcknowledgementsReferences