Development of high strength hot rolled low carbon copper-bearing steel containing nanometer sized carbides

M.P. Phaniraj a , 1, Young-Min Shin a , b, 1, Joonho Lee b , Nam Hoon

Goo c , Dong-Ik Kim a , Jin-Yoo Suh a , Woo-Sang Jung a , jae-Hyeok

Shim a , In-Suk Choi a , 

AbstractA low carbon ferritic steel was alloyed with Ti, Mo and Cu with the intention of achieving greater increment in strength by multiple precipitate strengthening. The steel is hot rolled and subjected to interrupted cooling to enable precipitation of Ti–Mo carbides and copper. Thermodynamic calculations were carried out to determine equilibrium phase fractions at different temperatures. Microstructure characterization using transmission electron microscopy and composition analysis revealed that the steel contains ~5 nm size precipitates of (Ti,Mo)C. Precipitation kinetics calculations using MatCalc software showed that mainly body centered cubic copper precipitates of size < 5nm form under the cooling conditions in the present study. The steel has the high tensile strength of 853 MPa and good ductility. The yield strength increases by 420 MPa, which is more than that achieved in hot rolled low carbon ferritic steels with only copper precipitates or only carbide precipitates. The precipitation and strengthening contribution of copper and (Ti,Mo)C precipitates and their effect on the work hardening behavior is discussed.

Keywords HSLA steel;  Copper;  Nano-sized carbides;  Interrupted cooling;  Numerical simulation

1. IntroductionThe long-standing goal in research on high strength steels has been to increase the strength while maintaining the ductility, toughness, weldability and cost effectiveness. In low carbon ferritic steels the strength–ductility combination is achieved by microalloying with elements such as Nb, V or Ti in the mill processed condition without the requirement of an additional heat treatment. The microalloying elements form nanocrystalline carbides/carbonitrides which strengthen the ferrite matrix by grain size refinement and precipitation strengthening. The conventional microalloyed steels typically possess yield strengths in the range 450–550 MPa [1] and [2], however significant improvements in strength

can be achieved by a combination of alloy design and thermo mechanical processing followed by controlled cooling.Recently, Funakawa et al. [3] reported an increase in yield strength of 300 MPa in the hot rolled low carbon steel micro-alloyed with Ti and Mo in equiatomic concentration. The hot rolled sheet was air cooled to 620 °C where phase transformation (γ→α) was accompanied by precipitation of alloy carbides. The high strength of the steel was due to ferrite grain size refinement (~3 μm) and precipitation of nanometer size (Ti,Mo)C carbides (~3 nm). Chen et al. [4] compared carbide size and hardness of three continuously cooled ferritic steels containing the microalloying elements Ti, Ti–Mo and Ti–Nb respectively. They reported that the Ti–Mo steel had the finest carbide size and highest hardness. The relatively lower change in carbide size and hardness measurements after different cooling rates lead them to conclude that the carbide formed in the Ti–Mo steel was relatively thermally stable.Copper addition in amounts of 1–2 wt% has also been used for precipitation strengthening in low carbon structural steels [5], [6], [7], [8] and [9]. Commercially available copper containing low carbon steels such as ASTM A710 or HSLA80, also contain microalloying elements such as Nb and have yield strength between 450 and 520 MPa [10] and [11] in the as-rolled and air-cooled condition. The higher strength of these steels is because of the formation of nanosized copper precipitates in ferrite and grain refinement brought out by microalloy carbides in prior austenite. Misra et al. [12] studied the effect of copper addition on the mechanical properties of hot rolled V–Nb microalloyed steels. They reported that increasing the copper content from 0.22 to 0.63 wt% increased the yield strength by ~25 MPa in the as rolled condition. In order to increase the strength further they microalloyed the Cu–Nb–V steel with Ti and Mo. The yield and tensile strength improved to 554 MPa and 658 MPa, respectively after aging the as-rolled specimen. The improvement in strength was due to ferrite grain refinement (2–10 μm), and nano sized precipitates of copper and the carbides: NbC, VC, Mo2C, Ti(Nb)C. The % elongation decreased with increase in strength. Their results show that multiple carbide precipitates may not guarantee the yield strength superior to that obtained in steels with fewer type of carbide precipitates [3] and [13].In the present study the low carbon steel is alloyed with both copper, and titanium and molybdenum with the intention of synergizing their beneficial effect on mechanical properties. The structure and properties of Ti–Mo and Cu–Ti–Mo microalloyed steels were investigated to determine the contributions to strength by Ti–Mo carbides individually and together with copper precipitates. The thermo-kinetic software MatCalc has been used to calculate evolution of copper precipitates in Fe–1.4% Cu alloy [14] and in austenitic heat resistant steels by this paper׳s co-authors [15]. In

the present study MatCalc was used to calculate the phase evolution and size distribution of copper precipitates.

2. Experimental and computational detailsThe compositions of steels designated as CMn, TiMo and CuTiMo are given in Table 1. CMn is the base low carbon steel without any precipitate forming elements. TiMo and CuTiMo are the CMn steel microalloyed with titanium and molybdenum, and microalloyed with copper, titanium and molybdenum respectively. The alloys were prepared using vacuum induction melting into 5.7 kg ingots. The ingots were forged and rolled at 1100 °C into 15 mm thick slabs. The slabs were then heated to 1250 °C and held for 30 min to dissolve any precipitates and then rolled to 75% reduction at 900 °C. After rolling the specimens were first air cooled to 650 °C and held for 5 min, followed by air cooling to 500 °C where it was held for 60 min and then furnace cooled to room temperature. The precipitates viz. (Ti,Mo)C carbide and Cu rich phase are expected to form at 650 °C [3], [4] and [16] and 500 °C [17] and [18] respectively.

Table 1.Chemical composition (wt%).Steel C Mn Si Mo Ti Cu Al FeCMn 0.0

71.47

0.32

– – – 0.04

Bal.

TiMo 0.07

1.34

0.32

0.20

0.09

– 0.04

Bal.

CuTiMo

0.07

1.53

0.34

0.21

0.09

1.17

0.04

Bal.

Table optionsThe specimens for the tension test were prepared from the rolled plates along the rolling direction according to ASTM standard E08-M with gage length of 25 mm, gage width of 6 mm and thickness of 2 mm. The tests were carried out at the constant crosshead speed of 1 mm/min. The tension test experiments were conducted on two specimens for each composition. The microstructure was characterized using scanning electron microscope (SEM), the 200 kV Tecnai 20 transmission electron microscope (TEM) and energy dispersive spectroscopy (EDS). Focused Ion Beam (FIB) technique was used to make thin foil specimens for characterization of precipitates in TEM-EDS. The precipitates were also extracted by dissolution in the electrolyte solution consisting of 4% tetramethylammonium chloride, 10% acetone and 86% methyl alcohol. The grain size was determined from SEM micrographs with help of an image analyzer using the linear intercept method. The precipitate size and distribution were determined from measurements in TEM micrographs on up to 400 extracted particles.The equilibrium phase fractions at different temperatures were calculated using Thermo-Calc [19] and [20]. The data set of parameters for the thermodynamic models describing the Fe–C–Mn–

Si–Mo–Ti–Cu system used in Themo-Calc was from the TCFE database incorporated in the software. The phases included in the calculations were austenite, ferrite, cementite, fcc-copper, and MC (M=Ti,Mo) carbide. A thermo-kinetic software MatCalc (version 5.30) [21], [22] and [23] was used to determine the precipitation kinetics and size distribution of Cu precipitates and MC carbides. In MatCalc the microstructure evolution is calculated based on the classical nucleation theory. The evolution equations for the radius and composition of each precipitate are derived from the thermodynamic extremum principle. The thermodynamic and kinetic data required for the simulation are calculated from the MatCalc database ‘mc_steel’, version 1.18, and the MatCalc mobility database ‘mc_sample_fe’, version 1.007. The following assumptions were made for the simulation: (i) average grain size and dislocation density of austenite is 50 μm and 1012 m−2, (ii) average grain size and dislocation density of ferrite is 15 μm and 1013 m−2, (iii) nucleation of bcc Cu and MC precipitates occurs on grain boundaries and dislocations, (iv) the austenite–ferrite transformation occurs at 650 °C and (v) the bcc Cu precipitates transform into fcc Cu precipitates when they grow in a range between 3–5 nm.

3. Results

3.1. Phase equilibriaThe equilibrium phase fractions calculated using ThermoCalc are plotted as a function of temperature in Fig. 1. In all the steels cementite precipitates below ~700 °C and precipitation of face centered cubic (fcc) copper begins at around 750 °C. In steels containing Ti and Mo MC carbide (NaCl structure) precipitates, consisting mainly of TiC, start to form around 1150 °C. In these steels the volume fraction of cementite is relatively lower than that in CMn steel. The lower cementite content is expected because some carbon is used up in the formation of MC type precipitates.

Fig. 1. Calculated equilibrium phase fractions plotted as a function of temperature in (a) CMn steel, (b) TiMo steel and (c) CuTiMo steel.

Figure options

3.2. MicrostructureThe microstructure in all the steels consists mainly of polygonal ferrite (Fig. 2). In the CMn steel cementite exists in pearlite form whereas in TiMo and CuTiMo steels pearlite colonies are less common. The latter is consistent with the calculations of equilibrium phase fractions (Fig. 1) where it was noted that the lower fraction of cementite was because some carbon was used in the formation of MC carbides. The ferrite grain size of CMn steel is 16 μm whereas both TiMo and CuTiMo steels have relatively finer grain size of 12 μm.

Fig. 2. SEM micrographs of the steels after hot rolling and cooling: (a) CMn steel, (b) TiMo steel and (C) CuTiMo steel.

Figure optionsFig. 3a is a TEM micrograph of the extracted nanometer size carbide particles from TiMo steel. The electron diffraction pattern from the precipitates shows that they have a typical face centered cubic structure (Fig. 3b). The lattice parameter calculated based on the diffraction ring diameter is 0.430 nm which is similar to that of TiC (0.433 nm). The EDS spectrum shows that MC precipitates contains Ti and Mo as main components (Fig. 3c). The copper peak (at ~8 eV) in the EDS spectrum is from the Cu grid that supports the carbon extraction replica. The composition analyses showed that MC precipitates contain molybdenum and titanium in the atomic concentration ratio between 0.2 and 0.4. This indicates that the precipitates are indeed (Ti,Mo)C precipitates. The precipitates have size in the range from 2 nm to 7 nm (Fig. 3d).

Fig. 3. (a) TEM micrograph showing precipitates extracted from TiMo steel, the corresponding, (b) fcc diffraction pattern, (c) EDS Spectrum from the precipitates, and (d) distribution of carbide size.

Figure options

Fig. 4 shows the diffraction pattern and bright field and dark field images from the foil specimen of the CuTiMo steel. The diffraction rings for (110)-αFe and for (111) and (200) from TiC are overlaid on the diffraction pattern. The diffraction pattern has been inverted i.e. the diffraction spots are now black and the background is white-grey to enable clearer identification of spots from carbide. The dark field image for the spot marked A in the diffraction pattern shows the size of the carbide particles. The particle size of the carbides in the present study was measured, like in the TiMo steel, from transmission electron micrographs of extracted precipitates and found to be in the range from 2 nm to 8 nm (Fig. 4d). Composition analysis showed that the Mo/Ti atomic concentration ratio, like in the TiMo steel, was in the range 0.2–0.4.

Fig. 4. (a) Bright field and (b) dark field image corresponding to the spot marked A in the diffraction pattern (c) from the CuTiMo steel and (d) size distribution of extracted precipitates.

Figure options

In the CuTiMo steels the copper precipitates could not be identified in the electron microscope. However, simulations using MatCalc for the processing temperature–time conditions employed in the present study showed that bcc-copper precipitation initiated in the ferrite matrix and reached saturation (Fig. 5a) and fcc-copper precipitates have begun to form. The precipitates are mainly body centered cubic and their size is largely between 2.4 nm and 3 nm (Fig. 5b). A small fraction of precipitates having size larger than 3 nm have transformed to the face centered cubic structure. The ~3 nm precipitates are fully coherent with the matrix and cannot be observed in the TEM because of poor diffraction contrast [24] and [25]. However, in a parallel unpublished work, by the present authors, on a 1.7 wt% Cu steel under similar processing conditions fcc-copper precipitates of size 6–13 nm formed and could be easily identified.

Fig. 5. Kinetics of copper precipitation in CuTiMo (a) and precipitate size distribution (b).

Figure options

3.3. Mechanical propertiesThe strength of the CMn steel increases with microalloying additions significantly (Fig. 6a). The yield stress,σy, of the TiMo steel (642 MPa) is more than 2.4 times that of CMn steel ( Table 2). The CuTiMo steel has the highest yield (732 MPa) and tensile strength (853 MPa). The % elongation of both TiMo and CuTiMo steels is similar (16%) but lower than that of the CMn steel (33%).

Fig. 6. Engineering stress-elongation (a) and work hardening rate-reduced stress (σ−σy) curves (b) for CMn, TiMo and CuTiMo steels.

Figure optionsTable 2.Measured yield strength (YS), tensile strength(TS), % elongation and calculated precipitation strengthening contribution and increment in YS due to finer grain size (Δσg).

SteelYS (MPa)

TS (MPa)

% Elongation

σppt a (MPa)

Δσg (MPa)

CMn 267 401 33 – –TiMo 642 770 16 354 21CuTiMo

732 853 16.6 417 21

aσppt is the increment in yield strength due to precipitates.

Table optionsIn order to compare the effect of alloying additions on the work hardening rates, the engineering stress–strain curves were first transformed into true stress–true strain curves. The data between the yield stress and the maximum stress were then fitted to a polynomial and this polynomial was then differentiated with respect to strain. Fig. 6b shows the work hardening rate, dσ/dε plotted as a function of the reduced flow stress, σ−σy, for the three steels. Typically there is a high initial work hardening rate followed by a decrease as deformation proceeds. The initial work hardening rate increases from ~3100 MPa in the CMn steel to ~4100 MPa in the TiMo steel. The CuTiMo steel has an initial work hardening rate marginally higher than the TiMo steel, however, the slope of the work hardening rate curve is initially lower and in the later stages of deformation is similar to the TiMo steel.

4. Discussion

4.1. Microstructure

The solubility of titanium in ferrite is low thereby cooling the steels after hot rolling to 650 °C supersaturates ferrite and results in the precipitation of alloy carbides during subsequent cooling. The alloy carbide precipitates contained molybdenum and titanium in the Mo/Ti atomic concentration ratio of 0.2–0.4 indicating that they could have formed over range of temperatures. A thermodynamic equilibrium calculation was carried out to determine the Mo/Ti atomic ratio as a function of temperature. The Mo/Ti atomic concentration ratio decreased with increasing temperature and reached 0.2 at ~600 °C (Fig. 7). Therefore the precipitates observed in the present study have most likely formed below 600 °C. Chen et al. [4] have reported similar precipitation from supersaturated ferrite when they cooled (Ti,Mo)C forming steels at different cooling rates.

Fig. 7. The atomic concentration ratio Mo/Ti in MC carbide precipitates in TiMo steel, obtained from thermodynamic equilibrium calculations, plotted as a function of temperature.

Figure optionsThe Fe-base copper alloys [5], [24] and [26] on aging, usually in the temperature range 400–650 °C, initially form metastable bcc copper precipitates. As these precipitates grow they transform to the 9R structure [27]and then at relatively larger size to the stable fcc structure. The metastable bcc structure is maintained up to a size of 2.4–3 nm [18] and [28]. The peak hardness during aging is achieved when the precipitates have the bcc structure. The size of copper precipitates in the CuTiMo steel is largely between 2.4 and 3 nm. Thereby, the strengthening due to copper precipitates in the CuTiMo steel should be analogous to that of the steel in the peak aged condition. Indeed the significant improvement in yield strength after adding copper to the TiMo steel, calculated in the next section, is mainly due to the bcc copper precipitates.

4.2. Estimating precipitation strengthening contributions

The base composition of the steels is similar to CMn steel and the

microstructure consists mainly of ferrite with some pearlite. Thereby

the contribution to yield strength after alloying can be estimated by

deducting the yield strength of CMn steel. The contribution to yield

strength from precipitates can be determined after accounting for

solid solution strengthening and grain size strengthening using

Eq. (1).

equation(1)

Turn MathJaxon

where Δσg is the increment in the contribution because of the slightly finer grain size of TiMo and CuTiMo steels, when compared with CMn steel, calculated using 17.402 d−1/2 [3] and [29] and σCu is the contribution from solid solution strengthening by copper (38 MPa per wt% Cu [1]). The amount of copper remaining in solution after precipitation was calculated to be 0.7 wt%; this amount was used to calculateσCu in Eq. (1). The solubility of titanium and molybdenum in ferrite is low and is assumed to have negligible solid solution strengthening effect. The precipitation strengthening contributions calculated using Eq. (1) are given in Table 2. The increment in yield strength due to precipitation in TiMo steel is 347 MPa. This is comparable to the 340 MPa improvement in yield strength reported on low carbon steels alloyed with Ti and Mo and processed under different conditions [3]. The precipitation strengthening contribution in the CuTiMo steel is 420 MPa which is more than one and half times the yield strength of the CMn steel.

4.3. Work hardening

In a matrix with non-shearable precipitates the moving dislocations

bypass the particle forming a loop around the particle rather than

cut it. Thereby, the effective spacing between particles is effectively

reduced and the stress required to move another dislocation past

the particle increases according to the Ashby–Orowan

mechanism [30]. The dislocation loop exerts a back stress which

must be overcome for additional slip to take place. Consequently

the matrix strain hardens rapidly as more dislocation loops form

around the particle. The (Ti,Mo)C precipitates, formed in the present

study, are hard and the dislocations should by pass the particles,

form loops around them, and thereby give rise to the significantly

higher work hardening rate (Fig. 6). Indeed dislocation loops have

been observed in deformed (Ti,Mo)C forming steels [31].

Copper precipitates are elastically soft and can be sheared by the

moving dislocations [26]. Nakashima et al.[32] have reported that in

Fe–Cu alloys moving dislocations cut and pass through copper

particles of size up to 70 nm. Once the particles are sheared the

resistance to further dislocation motion decreases which should

result in the lowering of the macroscopic strain hardening

rate [33] and [34]. The increase in strength on copper addition and

marginal increase in the initial work hardening rate (Fig. 6), when

compared with the TiMo steel, are in accordance with the

hypothesis noted above. However, the macroscopic work hardening

rate becomes similar to the TiMo steel in the later stages of

deformation. This could be explained as follows. The copper

precipitates in the present study have a size close to the size at

which transformation to the 9R structure [27] occurs. It is possible

that during deformation bcc copper particles with size around 3 nm

transformed to the stronger 9R structure which increased the

hardening rate such that the slope of the work hardening rate curve

becomes similar to the TiMo steel in the later stages of deformation

(Fig. 6b). Similar arguments have been posed by Militzer et

al. [35] in their study where Fe–Cu alloys exhibited higher work

hardening rate when the size of copper precipitates was close to the

size at which transformation occurs. However, further study is

needed to validate the hypothesis of the dynamic transformation

and the associated change in the hardening rate.

4.4. Ductility

Ductile failure is initiated by nucleation of voids at hard second

phase particles such as carbides and nitrides in steels. At large

strains the dislocation structure around the particle imposes high

local stress that upon reaching a critical value will cause decohesion

of the particle/matrix interface thereby nucleating a void. Particles

as small as 5 nm have been found to nucleate voids [36]. The

relatively lower ductility of the TiMo steel when compared with the

CMn steel is thereby likely due to void nucleation at the

carbide/matrix interface. Typically, increasing the volume fraction of

hard second phase decreases the ductility[37] and [38]. However,

there is no decrease in the ductility of the CuTiMo steel where the

matrix has copper precipitates in addition to carbide particles. This

is because, as noted above, the moving dislocations cut and pass

through the relatively soft copper particles resulting in only a

marginal increase in work hardening rate. Consequently the

increase in the local stress is lower than that required for particle–

matrix decohesion and the ductility remains unaffected.

5. Summary and conclusions

Three low carbon steels: a plain low carbon steel, microalloyed with

Ti and Mo, and microalloyed with Cu, Ti and Mo were prepared and

subjected to hot rolling followed by interrupted cooling. ThermoCalc

was used to determine equilibrium phase fractions at different

temperatures. MatCalc was used to simulate the precipitation

kinetics and determine size distribution of copper precipitate. The

microstructure was characterized using SEM, TEM and EDS. The

specimens were tested in tension to determine mechanical

properties.

1.

Alloying with Cu, Ti and Mo increases the yield strength of the

low carbon ferritic steel from 267 MPa to 732 MPa and tensile

strength from 401 MPa to 853 MPa.

2.

Fcc (Ti,Mo)C carbides of ~5 nm diameter and bcc Cu-rich

precipitates of <5 nm diameter form in the ferrite matrix.

3.

The increment in yield strength due to the precipitation

strengthening effect of carbides in the TiMo steel is 347 MPa.

The alloy carbides are hard and moving dislocations by-pass,

form loops around them giving rise to the significantly higher

work hardening rate.

4.

The precipitation strengthening contribution of copper and

carbide precipitates jointly, in the CuTiMo steel, is 420 MPa.

This is more than the increments in yield strength that would

be achieved with only copper precipitates or only carbide

precipitates. Though the total fraction of precipitates in the

CuTiMo steel has increased the ductility is not adversely

affected because of the shearable copper precipitates.