Effect of Molybdenum and Cooling Regime onMicrostructural Heterogeneity in Bainitic Steel Wires
Marc Ackermann,* Bernard Resiak, Pascal Buessler, Bertrand Michaut,and Wolfgang Bleck
1. Introduction
Material and process development is usually an optimization taskin which economic concerns and the demands for improvedmechanical properties are to be linked. In this regard, thedevelopment of steels with bainitic microstructure seemsvery attractive. However, bainitic microstructures require a
particularly precise process control, whichis closely matched with the chemical com-position of the material. Especially, bainiticsteels containing metastable retainedaustenite represent a concept, providingan outstanding combination of strengthand ductility.[1–5] It has been shown thatthe fatigue behavior could be improvedby the transformation of retained austeniteto martensite which retards crack growthdue to local compression stresses at thecrack tip.[3,4] Silicon delays cementite pre-cipitation due to its insolubility in carbidesand as a consequence retained austenite isstabilized by local carbon enrichmentwhereas the matrix consists of bainiticferrite.[5,6] In addition to bainitic ferriteas the primary phase, retained austeniterepresents the secondary phase as the filmbetween sheaves or as a blocky structure.Metastable austenite partly transforms intomartensite during cooling due to a carbongradient toward the center.[7] These blockystructures commonly resemble islands andare therefore commonly termed martensite–
austenite (M–A) islands. Molybdenum in these steels was observedto decrease the size of M–A[5] and in addition prevent manganeseembrittlement on grain boundaries.[8] According to severalauthors, coarsening of these M–A constituents causes thedeterioration of toughness by providing crack initiation sitesdue to a high degree of distortion within these structures.[9–11]
So far, M–A constituents were analyzed by several microscopymethods, but quantitative analysis only focused on providing anaverage or maximum size of M–A constituents.[6,9] This studyaims at a precise description of the microstructure–propertyrelationship in bainitic steels, in particular, by elaborating distri-butions of morphological parameters (M–A size, area, perimeter,and density) of blocky structures for bainitic steels. The impact ofthe cooling schedule variations that are possible in wire hotrolling on the microstructure distribution functions is investi-gated and conclusions with respect to the microstructural hetero-geneity and the final mechanical properties will be drawn.
2. Experimental Section
Two microalloyed medium-carbon steel grades were used withdifferent molybdenum contents but otherwise with similarchemical compositions (Table 1). The silicon content of 1.2 wt%
M. Ackermann, Prof. W. BleckSteel InstituteRWTH Aachen UniversityIntzestr. 1, 52072 Aachen, GermanyE-mail: [email protected]
B. Resiak, P. Buessler, B. MichautResearch and Development Bars and WiresArcelorMittal MaizièresVoie Romaine – BP 30320, Maizières-Lès-Metz Cedex 57283, France
The ORCID identification number(s) for the author(s) of this articlecan be found under https://doi.org/10.1002/srin.201900663.
Hot-rolled wire is often further processed into complex components andtherefore has to meet high demands on mechanical properties. Above all, duringwire production, close control of the cooling parameters after hot rolling isrequired if strength and toughness must be set within narrow limits. Parameterstudies are conducted in laboratory investigations to adjust bainitic micro-structures, that consist of bainitic ferrite as the primary phase, whereas retainedaustenite films and martensite–austenite (M–A) constituents represent thesecondary phase. Hot deformation trials with subsequent continuous coolingare conducted in a thermomechanical treatment simulator using two micro-alloyed steels with 0.25% C, 2% Mn, 0.03% Nb, and 0.03% Ti with or withoutMo. The cooling parameters are set according to the process window of thecooling conveyor at the wire rod rolling mill. A microstructural analysis showsinhomogeneities in the appearance of the secondary phase depending on thecooling schedule. A quantitative analysis of the microstructural constituentsindicates that the distribution function of the morphological characteristics ofthe M–A constituents corresponds to the mechanical properties. Therefore,cooling cycle and chemical composition are adjusted precisely to adjust thebainitic microstructure and achieve the desired mechanical properties.
delays carbide precipitation, whereas manganese was added forsolid solution strengthening and control of the phase transforma-tion behavior, especially the critical cooling rate. Chromiumwas added to constrict the bainitic phase field,[12,13] whereasmolybdenum and boron shifted the ferrite/pearlite phase trans-formations to lower cooling rates.[14–16] Titanium was added as astrong nitride-forming element to protect boron from BN precip-itation to keep boron in solution. Dissolved boron will segregateat grain boundaries and thus improve grain boundary strengthand as a consequence increase toughness.[15] Thus, the additionsof molybdenum and boron prevent manganese from embrittle-ment.[17] Niobium microalloying is generally used for prioraustenite grain size refinement and may also contribute byprecipitation strengthening.
Initially, ingots of 140� 140mm2 were melted in a vacuuminduction furnace. The melts were forged to 60� 60mm2, afteran annealing treatment at 1200 �C. Samples of 20� 20� 65mm3
were cut from the forged ingot for hot deformation trials using athermomechanical treatment simulator (TTS). Preliminary testswere conducted in a dilatometer to reach the process windowof the cooling path of the 16–18mmwire rod. Finally, two coolingprograms were set according to the industrial cooling conveyorprocess parameter window (Figure 1). For both cooling profiles,
austenitization was set at 1200 �C for 10min with subsequenthot deformation at 900 �C with 0.3 and 10 s�1 as applied strainand strain rate in the final rolling step, respectively (Table 2).The first cooling profile (regime I) intended to produce ahomogeneous lower bainite microstructure, with a cooling rateof 5 K s�1 from deformation temperature to quench-stop temper-ature at 400 �C. Afterward, the slowest possible cooling of0.3 K s�1 was applied according to the industrial process window.The second cooling profile (regime II) aimed at a mixed bainite/martensite microstructure. The temperature profile showedcooling with 2 K s�1 from hot deformation temperature at 900to 500 �C followed by simulated air cooling of 1 K s�1. Afterconducting TTS trials, secondary samples were extracted formicrostructural analysis by scanning electron microscopy (SEM),for Charpy V-notch tests and for tensile tests. Samples for themicrostructural observation were ground and subsequently pol-ished to a surface finish of 1 μm. The microstructure was revealedusing Nital etching, and several SEM micrographs were takenalong the horizontal and the vertical axes of the cross section every3mm to track microstructural changes along the cross section.
In a first approach, several SEM images per position were usedfor measuring the diameter of blocky constituents in the micro-structure, as proposed in the literature.[10] During this analysis,different shapes of M–A constituents (round vs more complexshapes/elongated) were observed. The size or diameter forsuch complex shapes was ambiguous. Therefore, in a secondapproach, the same images were used to measure furthermorphological M–A features (area, perimeter, and density) toget more reliable data than using only the diameter. In addition,the distance between adjacent M–A islands was measured andstatistically evaluated in frequency plots (Figure 2). These plotsprovided information on the appearance of rather closely linkedor less networked M–A islands. In this way, the result revealedcritical network structures of M–A constituents which facilitatedcrack propagation and therefore indicated a loss in toughness.
The TRansformation Induced Plasticity (TRIP) effect was ana-lyzed by measuring retained austenite fractions from tensile testsamples via X-ray diffraction (XRD) using XSTRESS3000 G2Rinstrument with Cr Kα radiation. The samples were groundand polished in a longitudinal direction with subsequent electro-polishing in A2 electrolyte (40 V for 10 s). All samples weremeasured at defined distances from the fracture surface forretained austenite fractions.
3. Results
For steel 1 and steel 2, the cooling cycles are superimposed withthe corresponding continuous cooling transformation (CCT)diagrams in Figure 3 and 4, respectively. The two tested coolingregimes in the CCT diagrams pass through the bainite phasefield, while avoiding any ferrite/pearlite formation. The first
Table 1. Chemical compositions of laboratory steels in wt%.
Figure 1. Schematic representations of the heat treatment for steel 1and steel 2. TA, tA, TU, φ, φ, Ttrans, T1, and T2 represent austenitizationtemperature, time, deformation temperature, strain rate, strain, transitiontemperature, cooling rate 1, and cooling rate 2, respectively.
Table 2. Heat treatment parameters for regimes I and II. TA, tA, TU,φ, φ, Ttrans, T1, and T2 represent austenitization temperature, time,deformation temperature, strain rate, strain, transition temperature,cooling rate 1, and cooling rate 2, respectively.
cooling profile—regime I—(marked in blue) enters the lowerbainite phase, whereas regime II (dashed in red) leads to a phasetransformation already at higher temperatures. Surprisingly,Ms temperature of steel 1 increases below the bainitic phase field,whereas steel 2 shows a slight decrease, becoming more signifi-cant with prior ferrite/pearlite transformation.
It is observed that molybdenum delays the ferrite and pearlitephase transformation, as illustrated by the CCT diagrams forsteel 1. The bainite phase field of the molybdenum-containingsteel 1 appears to be more restricted in terms of the transforma-tion temperature area, resulting in a narrow process window toproduce a bainitic structure. The measured martensite starttemperature Ms for steel 1 of 358 �C is lower compared withsteel 2 with 368 �C. A calculation of Ms by the equation ofAndrews yields 380 �C for steel 1 and 381 �C for steel 2[10]
Ms ¼ 512� 453C� 16.9Niþ 15Cr� 9.5Moþ 217C2
� 71.5CMn� 67.6CCr(1)
For the cooling regimes with t8/5 times (cooling time from 800to 500 �C) of nearly 60min (290 HV10 Vickers hardness) and120min (233 HV10), a transformation-free gap evolves betweenpearlite and bainite. In Mo-free grade steel 2, the bainite phasefield is broader in terms of transformation temperature but nar-rower in terms of the cooling rate compared with the Mo-bearinggrade. A transformation-free area can be seen at t8/5 of around13min (271 HV10) and 26min (282HV10). In the CCT diagrams,Vickers hardness is indicated for each cooling curve. In general,higher hardness levels are obtained for the Mo-bearing steel 1.
After hot deformation with subsequent cooling in regime I,both steels reveal a rather homogeneous refined lath-shapedmicrostructure composed of M–A constituents in additionto the packets of bainitic laths (Figure 5a,b). In regime II,both steels obtain a rather nonhomogeneous microstructure(Figure 5c,d). For example, fine lath-shaped areas can be foundnext to coarse areas with a large gradient in laths’ length andthickness. An absence of molybdenum in steel 2 seems to further
M-A distance distribution in µm
Cou
nts
[-]
Figure 2. Exemplary distance measurement of M–A islands between adjacent neighbors. Critical M–A network structures are revealed by transformationof distance data into frequency plots.
Figure 3. CCT diagram of steel 1. The cooling regimes for a homogenouslower bainitic microstructure (regime I in blue) and inhomogeneousbainitic/martensitic microstructure (regime II dashed in red) are indicated.Numbers in circles represent the Vickers hardness in HV10 for thecorresponding cooling curve. t8/5 time provides the cooling time from800 to 500 �C for samples with bainite phase transformation.
Figure 4. CCT diagram of steel 2. The cooling regimes for a homogenouslower bainitic microstructure (regime I in blue) and inhomogeneousbainitic/martensitic microstructure (regime II dashed in red) are indicated.Numbers in circles represent the Vickers hardness in HV10 for thecorresponding cooling curve. t8/5 time provides the cooling time from800 to 500 �C for samples with bainite phase transformation.
increase the microstructural inhomogeneity in regime II. Thiscoincides with the observations in the dilatometer experiments.Acquired data of length change and temperature were trans-formed into phase fraction over the temperature using the leverrule. For this approach, applied tangents on the linear sections ofthe cooling curve provide phase transformation start and finishtemperatures. A comparison of the phase fraction developmentduring cooling according to regime II shows that steel 2generates an increased transformed volume fraction at highertemperatures compared with steel 1 (Figure 6).
The observed fracture surface after Charpy V-notch testingof the mixed bainitic/martensitic cooling regime II appearsintercrystalline as pronounced brittle fracture (Figure 7), whereasespecially for the molybdenum-containing steel 1, after coolingalong the lower bainite region, dominantly dimples can beobserved, representing a rather ductile behavior.
The microstructural distribution of the M–A islands was char-acterized by number, average size, standard deviation of size, andmaximum size, as shown in Table 3. The comparison ofsteel 1 and 2 indicates that the Mo-free steel 2 in regime II showsa larger standard deviation in size of 0.42 μm and in total, a largerscattering of blocky constituents, in addition to an enlargedmaximum size of 2.3 μm and mean size of 1.4 μm. The cumula-tive counts of perimeter (Figure 8) show a higher degree of inho-mogeneity after applying the cooling parameters of regime II.This tendency appears stronger (or more inhomogeneous) forsteel 2 with a significant increase in the standard deviation ofperimeter from 1.61 to 6.20 μm (Table 4). The maximum perim-eter increases for steel 2 from 11.6 to 27.8 μm in regime II. Themeasured area per M–A island reveals the same tendencies as forthe perimeter with an increased scatter of area (“SD” value of 1.1and 3.5 μm) and coarsening of the maximum area (“Max” valueof 5.6 and 18.5 μm) for cooling regime II (Table 5). The coarsen-ing effect by the upper bainite cooling variant is significant inboth steels especially for the results of M–A perimeter and area.The frequency plots of distances between nearest M–A neighbors(Figure 9) illustrate an accumulation of short distances for theMo-free steel 2, regardless of applied cooling. The observeddegree of short distances increases further for regime II.Mo-containing steel 1 in regime I seems to obtain the lowestconcentration of short distances among the tested samples.
The mechanical properties of samples tested in regime Iappear similar in terms of strength and elongation with only
Figure 5. Microstructures of a,c) steel 1 and b,d) steel 2 in cooling a,b) regime I and c,d) regime II.
200 300 400 500 6000,0
0,2
0,4
0,6
0,8
1,0
Pha
se F
ract
ion
Temperature in°C
Steel 2
Steel 1
Regime II
Figure 6. Phase transformation during cooling according to regime II ofsteel 1 and 2. The absence of molybdenum is seen with a shift of phasetransformation to higher temperatures for steel 2.
a rather small decrease in impact toughness for steel 2 inregime II (Table 6). In contrast, samples in regime II show alarger deviation. For example, the absence of molybdenum causes
in regime II a drop in tensile strength by 80MPa but slightlyincreases the yield strength of steel 2 compared with steel 1,resulting in an increased yield ratio. Total elongation almostdoubles without molybdenum for the same cooling profile.In contrast, the molybdenum addition in steel 1 causes a higherimpact energy. The calculated average of impact energy is basedon the measurements of 3–4 undersized samples of 55� 10�2.5mm3 per condition with a rather narrow scattering (Figure 10).
The T 00 temperature provides information on the maximum
formed banitic ferrite at a given temperature, depending on theoverall carbon content and carbon concentration in austenite.[18]
At this temperature, further growth of bainitic ferrite is thermo-mechanically impossible due to a lower free energy of bainitecompared with austenite (considering the stored energy in ferriteof 400 J mol�1).[19] The remaining volume fraction represents
Figure 7. Fracture surface after Charpy testing at room temperature of a,c) steel 1 and b,d) steel 2 at cooling a,b) regime I and c,d) regime II.
Table 3. Influence of cooling schedule on M–A size parameters for steel 1and steel 2 conditions (N, Mean, SD, and Max as number, average,standard deviation, and maximum of measured M–A size, respectively).
Size N [–] Mean [μm] SD [μm] Max [μm]
Steel 1 Regime I 52 0.6 0.14 1.0
Regime II 30 0.7 0.21 1.1
Steel 2 Regime I 33 0.6 0.22 1.2
Regime II 50 1.4 0.42 2.3
0 5 10 15 20 25 300
10
20
30
40
50
Cou
nts
[-]
M-A Perimeter in m
0 5 10 15 20 25 300
10
20
30
40
50
Cou
nts
[-]
M-A perimeter in m
(a) (b)Steel 1 Steel 2
Regime I
Regime II
Regime I
Regime II
Figure 8. Frequency plots of the M–A perimeter showing the effect of the cooling regime for a) steel 1 and b) steel 2.
retained austenite as an incomplete transformation product. Thecalculated curve shows a slight shift of T 0
0 (Figure 11) to lowercarbon for the Mo-bearing steel 1, indicating a higher degreeof incomplete transformation compared with steel 2.
The stability of retained austenite can be monitored bymeasuring austenite fraction before and after tensile deforma-tion by means of XRD. The XRD results of the tensile test sam-ples exhibit at 0% strain a high level for the retained austenitefraction between 7 and 10 vol%, with steel 2 showing slightlyhigher amounts (Figure 12). After deformation, the retainedaustenite fraction decreases for both steels depending on thedistance to the fracture surface—in other words, depending onthe amount of strain—approaching a complete transformation tomartensite next to the fracture surface. Thus, all tested steelsshow a TRIP effect after tensile testing at room temperature.The higher initial retained austenite fraction of steel 2 decreasesrapidly, leading to a lower fraction at 7 mm distance and
subsequently to a similar retained austenite response com-parable with steel 1 in regime II. For regime I, with a phase trans-formation mainly in the lower bainite area, the Mo-free gradesteel 2 surpasses the fraction of retained austenite of regime Iin steel 1. But in general, it has to be mentioned that the differ-ence lies within the range of the statistical error.
Table 4. Influence of cooling schedule on M–A perimeter for steel 1 andsteel 2 conditions (N, Mean, SD, and Max as number, average, standarddeviation, and maximum of measured M–A perimeter, respectively).
Perimeter N [–] Mean [μm] SD [μm] Max [μm]
Steel 1 Regime I 34 3.8 1.46 7.5
Regime II 27 6.5 3.87 21.3
Steel 2 Regime I 45 4.2 1.61 11.6
Regime II 42 11.4 6.20 27.8
Table 5. Influence of cooling schedule on M–A area for steel 1 and steel 2conditions (N, mean, SD, and Max as number, average, standarddeviation, and maximum of measured M–A area, respectively).
Area N [–] Mean [μm] SD [μm] Max [μm]
Steel 1 Regime I 34 0.6 0.4 1.8
Regime II 27 1.3 1.1 5.6
Steel 2 Regime I 45 0.7 0.5 3.4
Regime II 42 3.7 3.5 18.5
0 10 20 30 400
10
20
30
40
50
Cou
nts
[-]
M-A distance distribution in mM-A distance distribution in m0 10 20 30 40
0
10
20
30
40
50
Cou
nts
[-]
(b)(a)
Steel I
Steel 2
Steel I
Steel 2
Regime I Regime II
Figure 9. Influence of molybdenum on the M–A density via distance measurements of adjacent neighbors of M–A for a) regime I and b) regime II;Mo-bearing steel 1 and Mo-free steel 2 marked in blue and red, respectively.
Table 6. Mechanical properties of regime I and II for steel 1 and (Mo-free)steel 2. YS, UTS, UEl, Tel, and CIV as yield strength, ultimate tensilestrength, uniform elongation, total elongation, and Charpy impact value(2.5mm-sized samples), respectively.
YS[MPa]
UTS[MPa]
Yieldratio [–]
UEl[%]
TEl[%]
CIV[ J]
Steel 1 Regime I 925 1344 0.69 3.9 10.5 18
Regime II 921 1367 0.67 2.8 5.3 10
Steel 2 Regime I 965 1365 0.71 3.8 10.3 15
Regime II 931 1287 0.72 3.7 10.7 6
Regime I Regime II Regime I Regime II0
5
10
15
20
Impa
ct e
nerg
y in
J
Steel 1 Steel 2
Figure 10. Charpy impact energies tested at room temperature(undersized samples with 2.5 mm in thickness).
For both steels, a rather narrow size distribution of M–A constit-uents can be seen in regime I. M–A constituents become coarser(increased mean size and maximum size) and more scattered(increased standard deviation of size distribution) in regimeII, especially for the Mo-free steel 2. The size distribution interms of standard deviation of steel 1 appears similar for bothcooling regimes. In contrast, the perimeter and the area distri-bution of M–A constituents result in a larger standard deviationfor regime II of steel 1, indicating an increased inhomogeneity.This difference in contrast to the M–A diameter can be explainedby different observed shapes of the M–A constituents which are
usually not regular circles. In case of a circular M–A island, withan aspect ratio approaching a value of 1, excess carbon candiffuse from surrounding bainitic sheaves but only penetratesthe outer edge of M–A. The round-shaped M–A island containsa carbon gradient toward the center, leading to a lower stability ofretained austenite. In case of elongated or complex-shaped M–Aconstituents, carbon diffuses from surrounding bainitic sheaves,but as the diffusion distance is shorter for penetrating elongatedM–A, the carbon concentration is more homogeneous and prob-ably higher. For this example, the complex-shaped constituentrepresents a larger perimeter relative to the round-shaped M–Aisland for the same area. Thus, the (average) size or diameter ofM–A islands seems not to be solely meaningful. The additionalinformation on M–A perimeter (or area) provides therefore fur-ther insights into M–A morphology. It is further recommendedto identify networks of M–A islands as critical regions for facili-tated crack propagation. The presented method addresses thedistances of nearest neighbors of M–A constituents and in caseof an accumulation of short distances, network structures can beidentified, as it was for Mo-free steel 2, regime II. The coarseningof M–A constituents can be linked to the higher quasi-isothermalholding temperature of 500 �C after cooling from 900 �C, with2 K s�1 facilitating carbon partitioning from bainitic ferrite toadjacent austenite. In combination with a shorter time for bainiteformation within the bainitic phase field, larger retained austen-ite islands are formed, which partly transform into martensiteand finally produce in average coarser M–A islands.
In general, a drop in impact energy seems to correlate with ahigher degree of microstructural inhomogeneity, which isobserved for regime II, especially for steel 2. For example, adecreased maximum perimeter (and less scattering of the samedue to a reduced standard deviation) corresponds to an enhancedimpact energy of both steels (Figure 13b), whereas solely the sizeof M–A is not sufficient to explain the effect on impact energy, asno significant difference can be seen for regime I of steel 1(Figure 13a).
4.2. Effect of Molybdenum and Cooling Regime
The influence of molybdenum on the phase transformationbehavior results in a slight decrease in Ms temperature for thequenched samples of Mo-bearing steel 1. This tendency isexpected when the empirical formula by Andrews is applied,but the effect is rather insignificant. At lower cooling rates inthe CCT diagram (t8/5 time of 48–384 s), the austenite to bainitetransformation causes carbon diffusion into adjacent retainedaustenite. As a result, the remaining austenite becomes stabilizedand therefore causes a decrease inMs temperature. This tendencywas observed for steel 2. In contrast, steel 1 shows an increase inMs temperature subsequent to bainite transformation. Retrialof the dilatometer experiments suggests the same observation.It is assumed that molybdenum causes precipitation of MoCcarbides, depleting the remaining austenite with carbon and thus,Ms temperature increases. The same tendency ofmolybdenum onMs was already argued by Capdevila et al.[20] Ongoing analysisshould clarify the appearance and quantity of MoC carbides.
Furthermore, both steels show a different response to the twoinvestigated cooling regimes. The role of molybdenum appears
0,01 0,02 0,03 0,04
300
400
500
600
Tem
pera
ture
in °
C
Mol fraction C
T0'
Steel 2Steel 1
Figure 11. Calculated T 00 curve for steel 1 and steel 2. Steel 2 shows a slight
shift to a higher carbon concentration.
0 1 2 3 4 5 6 7 9 10 11 120
1
2
3
4
5
6
7
8
9
10
11
Ret
aine
d au
sten
ite fr
actio
n in
vol
.%
Distance from fracture surface in mm
0% s
trai
n
Steel 2
Steel 1
1mm 3mm 5mm 7mm
Figure 12. Evolution of retained austenite over distance from fracturesurface after tensile testing. Reference values are based on nondeformedsamples (0% strain).
in regime I as less significant, whereas in regime II, an absenceof molybdenum seems to pronounce both the effect of coarsen-ing and forming critical network M–A structures. Thus, molyb-denum has a refining effect on blocky constituents linked to lesscritically distributed blocky M–A.
The influence of cooling parameters on the mechanical prop-erties is more significant in regime II, in particular, visible in thedeterioration of impact energy for both steels. In steel 1, addition-ally, the total elongation decreases. This drop can be explained bya smaller amount of soft bainitic ferrite. For steel 2, no change intotal elongation was observed for the two cooling cycles, althoughthere exist some differences in the microstructure homogeneityfor the two cooling regimes. One difference in the microstruc-ture of steel 2 in regime I is regarding a higher amount of bainiticferrite. Thus, in contrast to steel 1, the total elongation can bemaintained for both regimes. The amount of bainitic ferritecan be adjusted by the overall carbon content and substitutionalalloying elements. According to Bhadeshia and Edmonds, anincrease in volume fraction Vαb of bainitic ferrite correlates withthe reduction of overall carbon content x, assuming constantcarbon concentration in bainitic ferrite xαb and in retainedaustenite xT0
(at temperature T0)[2,5]
Vαb ¼ðxT0
� xÞðxT0
� xαbÞ(2)
In addition to the carbon content, the resulting T0 curve for atemperature range of bainitic phase transformation can be shiftedby the substitutional elements. A reduction in substitutionalelements causes a shift to higher carbon concentrations[18] andthus, a higher amount of bainitic ferrite can be expected. Thiscan be seen by the T 0
0 curve (Figure 11) for steel 2 with a highercarbon concentration compared with Mo-containing steel 1 at thesame temperature. This results in a higher amount of bainiticferrite of steel 2 in accordance with the SEM observations.
The impact of molybdenum can be also seen on impact tough-ness. The addition of 0.08 wt% Mo increases the Charpy impactenergy, independent of the cooling regime. According to theliterature, additional molybdenum causes a higher transitiontemperature from plate to sheave morphology.[21] Thus, a higherdegree of bainitic sheaves in the investigated microstructures of
steel 1 with favorable high-angle misorientations would be able toabsorb a higher impact energy.
Another supporting effect of molybdenum on impact tough-ness can be argued with the combined use with boron. In steel 1,the combination of MoþB seems to avoid grain boundaryembrittlement, more effective than with the sole use of B.This tendency can be observed from the fracture surface ofsteel 1 and steel 2, where especially steel 1 can be linked to a lackof intercrystalline fracture and therefore to a higher impactenergy compared with steel 2.
4.3. TRIP Effect
The XRD measurements of tensile test specimens indicatea similar behavior of both steels. Close to the fracture surface,the retained austenite is almost entirely consumed. In the less-strained region of the specimens, differences exist in retainedaustenite stability between the Mo-bearing and the Mo-free steel.The Mo-alloyed steel 1 obtains a lower retained austenite fractionbefore deformation but maintains a higher fraction duringstraining (Figure 12). In contrast, steel 2 consumes some ofthe less stable retained austenite. This can be observed at thedecrease in retained austenite from the undeformed conditiontoward the fraction at 7 mm distance. In steel 2, the transformedphase fraction from dilatometry data suggests a higher amountof transformation already at higher temperatures compared withMo-bearing steel 1, especially for regime II (Figure 6). Fromboth, the dilatometer data and the SEM micrographs, it isbelieved that the higher retained austenite fraction of steel 2in regime II can be explained by the presence of granular bainitewith a higher degree of blocky austenite with a lower stability.Therefore, the impact of molybdenum on the retained austenitestability is more significant for regime II.
5. Conclusions
The influence of two different cooling regimes on the bainiticmorphology of molybdenum-containing and a molybdenum-freesteel was investigated. M–A constituents were measured andanalyzed in terms of size, area, and perimeter distribution,
Max M-A size in m mMax M-A perimeter in 0 1 2 3
0
5
10
15
20
0 5 10 15 20 25 30
Impa
ct e
nerg
y in
J
0
5
10
15
20
Impa
ct e
nerg
y in
J
(a) (b)Steel 1Steel 2
+Mo Steel 1Steel 2
Regime II
Regime I
+Mo
Figure 13. Effect of a) maximumM–A size and b) maximumM–A perimeter on impact energy. The M–A size is not sufficient to understand the influenceon impact energy for steel 1.
in addition to the M–A density. The findings can be summarizedas follows: 1) for the investigated steels, a cooling regimecrossing the upper bainite regime causes a high degree of micro-structural inhomogeneity in contrast to a quasi-isothermalcooling profile within the lower bainite phase field. 2) A highermicrostructural inhomogeneity causes lower ductility. Especially,the morphological distribution of the M–A constituents isrelevant for toughness. 3) In addition to the average size of blockyconstituents, the M–A perimeter, their area, and their dis-tribution provide more insights into the homogeneity of themicrostructure. Together with the density of the M–A constitu-ents, this information reveals critical network structures.4) Molybdenum causes a refinement and a narrower sizedistribution of M–A constituents. Moreover, molybdenum hasa stronger impact on properties when upper bainite is formedduring cooling. The molybdenum effect on strength and ductilityis rather insignificant for isothermal lower bainite formation.5) Molybdenum can be used to refine M–A constituents, improv-ing the impact energy. 6) For all steels, a TRIP effect wasobserved. The retained austenite stability can be controlled bythe molybdenum content and the cooling regime.
AcknowledgementsThe work was done in the framework of a collaboration project withArcelorMittal Maizières, Research and Development Bars and Wires, aspart of the knowledge building program at ArcelorMittal.
Conflict of InterestThe authors declare no conflict of interest.
Received: December 17, 2019Revised: February 27, 2020
Published online:
[1] V. T. T. Miihkinen, D. V. Edmonds, Mater. Sci. Technol. 1987, 3, 432.[2] H. K. D. H. Bhadeshia, D. V. Edmonds, Met. Sci. 1983, 17, 411.[3] F. G. Caballero, S. Allain, J. Cornide, J. D. Puerta Velásquez, C. Garcia-
Mateo, M. K. Miller, Mater. Des. 2013, 49, 667.[4] T. B. Hilditch, I. B. Timokhina, L. T. Robertson, E. V. Pereloma,
P. D. Hodgson, Metall. Mater. Trans. A 2009, 40, 342.[5] H. K. D. H. Bhadeshia, D. V. Edmonds, Met. Sci. 1983,
17, 420.[6] P. Jacques, E. Girault, T. Catlin, N. Geerlofs, T. Kop, S. Van Der
Zwaag, F. Delannay, Mater. Sci. Eng. A 1999, 273–275, 475.[7] C. Hofer, F. Winkelhofer, H. Clemens, S. Primig, Mater. Sci. Eng. A
2016, 664, 236.[8] P. Dumoulin, M. Guttmann, M. Foucault, M. Palmier, M. Wayman,
M. Biscondi, Met. Sci. 1980, 14, 1.[9] A. Lambert, J. Drillet, A. F. Gourgues, T. Sturel, A. Pineau, Sci. Technol.
Weld. Joining 2013, 5, 168.[10] F. G. Caballero, H. Roelofs, S. Hasler, C. Capdevila, J. Chao,
J. Cornide, C. Garcia-Mateo, Mater. Sci. Technol. 2013, 28, 95.[11] X. Li, Y. Fan, X. Ma, S. V. Subramanian, C. Shang, Mater. Des. 2015,
67, 457.[12] C. Keul, V. Wirths, W. Bleck, Archiv. Civil Mech. Eng. 2012, 12, 119.[13] A. Samoilov, Y. Titovets, N. Y. Zolotorevsky, G. Hribernig, A. Pichler,
Mater. Sci. Forum 2005, 500–501, 311.[14] B. C. de Cooman, Curr. Opin. Solid State Mater. Sci. 2004, 8, 285.[15] K. Zhu, C. Oberbillig, C. Musik, D. Loison, T. Iung, Mater. Sci. Eng. A
2011, 528, 4222.[16] R. A. Grange, T. M. Garvey, Trans. Am. Soc. Met. 1946, 37, 131.[17] I. Tsukatani, S. I. Hashimoto, T. Inoue, ISIJ Int. 1991, 31, 992.[18] F. G. Caballero, M. J. Santofimia, C. Capdevila, C. García-Mateo,
C. García de Andrés, ISIJ Int. 2006, 46, 1479.[19] H. K. D. H. Bhadeshia, Acta Metall. 1981, 29, 1117.[20] C. Capdevila, F. G. Caballero, C. García de Andrés, ISIJ Int. 2002,
42, 894.[21] C. García de Andrés, C. Capdevila, I. Madariaga, I. Gutiérrez, Scr.
