Tensile properties and bake hardening response of dual phase steels with varied martensite volume fraction

Journal Pre-proof Tensile properties and bake hardening response of dual phase steels with varied martensite volume fraction Mohamed Soliman, Heinz Palkowski PII:

S0921-5093(20)30132-5

DOI:

https://doi.org/10.1016/j.msea.2020.139044

Reference:

MSA 139044

To appear in:

Materials Science & Engineering A

Received Date: 23 October 2019 Revised Date:

31 January 2020

Accepted Date: 2 February 2020

Please cite this article as: M. Soliman, H. Palkowski, Tensile properties and bake hardening response of dual phase steels with varied martensite volume fraction, Materials Science & Engineering A (2020), doi: https://doi.org/10.1016/j.msea.2020.139044. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

Tensile properties and bake hardening response of dual phase steels with varied martensite volume fraction Mohamed Solimana, 1 and Heinz Palkowskia, 2 a

Institute of Metallurgy, Clausthal University of Technology, Robert-Koch-Straße 42, 38678

Clausthal-Zellerfeld, Germany 1

corresponding author: [email protected]

2

[email protected]

Abstract A study has been made of the tensile properties on martensite plus ferrite dual phase (DP) steels. Tensile samples machined out of industrially produced DP steel sheet are quenched from the austenite plus ferrite phase field, so as to give a series of microstructures with varying martensite volume fraction (MVF) from 0.17 to 1. The strength of the produced DP steels is found to linearly increase with increasing the MVF, whereas their ductility showed a nonlinear decrease. Tensile samples of the produced DP-steels were taken to study the bake hardening (BH) response. The samples were aged at 100, 170 or 220 °C after applying pre-straining of 0, 2% or 5%. The change of BH-response with the time is found to go through two stages, an initial small increase in yield strength, followed by a larger strength increase at longer baking times. Prolonging the aging time results in decreasing the BH-response achieved in the second-stage (over-aging). The BH-response in the first stage is improved by increased pre-straining. This trend implies that Cottrell atmosphere formation is highly influenced by pre-straining. However, during the second stage the effect is reversed, with the non-pre-strained condition recording higher BH-responses up to 180 MPa. It is found also that changing the aging temperature does not change the characteristics of the aging-phenomenon, but it affects its

1

rate. Saturation of increasing the BH-response with increasing MVF in the second strengthening stage is observed at MVF-range of 0.22-0.39. Keywords: dual phase steel; bake hardening; martensite volume fraction; tensile properties

1. Introduction It is an ongoing challenge for the automotive industry to create stronger and lighter components from sheet metal. A massive effort has been done to incorporate and develop high-strength steel materials. Dual phase (DP) steels are very attractive for the automotive industry owing to their combination of high strength coupled with good formability and excellent weldability. That, in addition to their good bake hardening (BH) response enable their use as an option for light weighting, increasing fuel efficiency [1]. The properties of DP steels rely on a microstructure of ductile ferrite with a reinforcing hard phase comprising of 10-40 vol.% of the microstructure. The hard-reinforcing phase is commonly martensite; however, it is not uncommon for amounts of retained austenite or bainite to likewise exist in the microstructure [2, 3]. Bhadeshia and Edmonds stated that the effects of martensite and bainite on the strength of DP steels are described by a common curve [4]. Other authors suggest that hardening by bainite is substantially weaker than that by martensite [5]. The martensite volume fraction (MVF) is a dominant factor in controlling strength and ductility of DP steels [6, 7, 8]. A linear relationship between the tensile strength of DP steel and the MVF was noted by many authors [7, 9, 10, 11]. Serious debates continue as to whether this relationship should depend on the properties of martensite or ferrite matrix [12]. A high BH-response is another interesting aspect of DP steels, enabling them to retain high formability throughout part production, while improving strength during the paint baking stage [3, 13].

2

The mechanism of bake hardening has been reported in literature to proceed through two stages. The first stage is the diffusion of carbon atoms to dislocations forming what is referred to as Cottrell atmospheres, while the second stage is the precipitation of carbides at dislocations. Both stages lead to the pinning of dislocations and hence strengthening of the bake hardened steel [3, 14, 15]. This study aims at investigating the effect of MVF on the tensile properties and bake hardening behavior of dual phase steel. The investigation was performed on an industrialproduced DP steel grade for the purpose of carrying out the study on a raw material with a typical chemical composition of the DP steels. In this study dual phase steels with varied MVF were produced out of this industrial grade by heat treating in a salt bath at different annealing temperatures. The annealing temperatures were selected following thermodynamic calculations and dilatometric investigations. Special care was taken to suppress the decarburization of the tensile samples. In this study, bainite transformation was excluded in the produced DP steel grades via applying extremely high cooling rate during quenching from the annealing temperature. Obtaining an uncontrolled mixture of martensite and bainite as a hard phase in the DP steel will significantly affect the dependence of the mechanical-behavior and the BH-response on the volume fraction of the hard phase and therefore would guide to misleading deductions. Therefore, suppressing the bainite transformation was fundamental to accomplish the aim of this study. In order to simulate the bake hardening (paint baking), strain aging was applied on samples with selected MVFs. The pre-straining is applied to simulate the deformation process, which takes place on the steel sheets prior to paint baking process. Deforming the steel results in increasing the dislocation density and therefore influences the BH-kinetics.

3

2. Experimental Procedure 2.1 As-received material and heat treatment preparations A 2-mm industrially produced hot rolled DP sheet with a nominal ultimate tensile strength of 600 MPa (DPW600) was studied. Table 1 gives the chemical composition of the material. The chemical analysis was performed using glow discharge spectrometry (GDS). The combustion method was used for accurate measurement of C and N. The C and N contents given in Table 1 rely on the latter measurement method. Table 1 Chemical composition of DPW600 (wt.%). C

N

Mn

Cr

Al

Si

S

P

0.075

0.007

1.02

0.433

0.031

0.059

0.005

0.020

2.2 Characterisation of the intercritical region and the heat treatment Throughout this study, the DPW600 was investigated to provide an insight of the effect of the martensite volume fraction (MVF) on the mechanical- and BH behavior. The approach of selecting the intercritical annealing temperatures took place in three consecutive steps; dilatometry, thermodynamic calculations and salt bath heat treatment. 2.2.1 Dilatometry The intercritical region was characterized using dilatometric measurements. This is achieved by slow heating in the intercritical region to trace the equilibrium points. The dilatometric measurements were conducted on a dilatometer Bähr “DIL 805A/D”, which performs a resolution of 0.05mm/0.05 K, using samples with 2 mm thick, 5 mm width and 10 mm length with the length in the direction of the dilatometric measurement. The test specimens had been degreased with an acetone solvent. Sheathed type S “Pt/Pt–10% Rh” thermocouples with a nominal diameter of 0.1 mm have been individually spot welded to the specimen’s surface in central position to monitor the temperature. The thermal cycles had been performed under 4

vacuum of 5×10-4 mbar. Helium was used for cooling. The dilatometric curves had been recorded along the thermal cycle with the help of a computer-data acquisition system. Detailed experimental procedure and device description are given elsewhere [16, 17]. Defining the intercritical region applying the dilatometric measurement will be discussed later. 2.2.2 Thermodynamic calculations Thermodynamic calculations were performed on the investigated material using Thermo-Calc employing the data base TCFE8. This software enables calculating phase equilibria, and then plot the results as temperature-phase fraction property diagram. 2.2.3 Salt bath heat treatment Based on the calculations of ThermoCalc and probed thru dilatometric investigation, different temperatures corresponding to different MVF in the DP-steel had been selected. Six different MVFs were produced by varying the intercritical annealing temperature (Ti). Additionally, the full martensitic microstructure was also produced. After annealing the samples, they were quenched in brine (10 % NaCl aqueous solution) to room temperature. A samples holder, shown in Fig. 1a., was fabricated to increase the consistency of the thermal cycle being applied on the tensile samples. This would give rise to the reproducibility of the tensile properties results as well. For monitoring the thermal-cycle, a 1 mm sheathed thermocouple of type K was inserted in a blind-hole bored perpendicular to the thickness in a middle-position of a dummy-sample and connected to a data-logger. The samples were austenitised in a Durferrit GS 540 salt bath, which has a working range of 600 to 950 °C. The molten salt tends to enrich itself with traces of oxygen compounds which cause surface decarburization. Decarburization was prevented by avoiding the reaction between the salt bath and the specimens by adding the inertor R2 to the salt. The inertor R2 is a barium-based powder. This substance has a high affinity to oxygen; therefore, it bounds the free oxygen of the salt-bath and settles to the crucible-ground. 5

The salt bath treatment followed the procedure: 1- Melting the salt by heating to the desired annealing temperature. 2- Dusting R2 into the bath, the added quantity was about 300 g per 10 l of molten salt. 3- Checking the effectivity of the inertor using the foil method. In this method a steel foil having 1% C and a thickness of 50 µm is immersed into the bath, removed after 30 min and then quenched. 4- After quenching the foil is bent between the fingers. The fragmentation of the foil means that the bath is neutral. Accordingly, the samples are inserted in the salt bath for the heat treatment (Fig. 1b). Otherwise, in case of bendable foil, more R2 must be added and the foil test repeated [18]. The heat treatments for the different annealing temperatures were carried out on the machined tensile samples for 20 min.

Dummy sample

Thermocouple

b

Thermocouple

a

Fig. 1 Tensile samples (a) inserted in the sample-holder and (b) immersed in a molten salt bath.

6

2.3 Microstructure characterization Light optical microscopy (LOM) and scanning electron microscopy (SEM) were performed on sections cut perpendicular to the width-direction of the samples. A slight tempering at 200 °C for 2 h before mechanical preparation and etching is adopted throughout this study. The tempering treatment is essential to enable distinguishing the martensite substructure and differentiating it within ferrite as will be seen later. After the slight tempering, the samples were mounted, ground, using standard abrasive grinding papers ranging from course 180 grit to fine 1200 grit, and finally polished until 0.3 µm OPS. The microstructures were examined after etching with 3% nital etchant. Some samples were etched with LePera’s etchant as well. The LePera’s etchant composed of two parts: 1 grams of sodium metabisulfite + 100 ml distilled water and 4 grams of picric acid + 100 ml ethyl alcohol. The two components are kept separate until use when they were mixed together in equal parts by volume [19]. Stereological measurements were carried out to evaluate MVF using the manual point count method according to ASTM E562 [20]. For each condition, the measurements were performed on ten fields using a square-grid with 144 points. The ferrite grain size was measured applying the mean linear intercept method according to ASTM E112 [21]. For each condition, 10 fields, each one was covered with 10 lines are considered.

2.4 Tensile testing The materials were machined to the required tensile specimen size and geometry prior to heat treatment to avoid possible strain ageing during machining. Standard subsize tensile specimens with a 6.4 mm width and a 25.4 mm gage length were machined in accordance to ASTM standard E8-13 [22]. The specimens were machined longitudinal to the rolling direction. A 25-ton computerized universal testing machine equipped with a video extensometer was used for conducting the tensile test until fracture as well as pre-straining the tensile samples. 7

For accurate identification of the yield points, different loading rates were applied as follows [23]: 1- 10 MPa/s during the elastic region. 2- 0.0115 mm/s during the yielding stage. 3- 0.0067 s-1 after yielding (plastic region) until fracture. Please note that the accurate identification of the yield point is crucial for a precise measurement of the BH-response.

2.5 Bake hardening Three different levels of pre-strain (0, 2 and 5 %) were applied on the DP steels. Up to eleven aging times, ranging from 2 min up to 20, 000 min., were applied at ageing-temperatures of 100, 170 and 220 °C. The investigated time-temperature combinations of the BH-parameters do not only cover the industrial paint baking parameters, but they additionally cover parameters out-of-industrial-scope. The pursue of covering a wide range of parameters was to contribute in understanding the strain aging phenomenon, by and large, in the DP steels. The heating time to the aging temperature in air circulated furnace was measured to be about 2 min. The significance of this period with respect to the aging time is high when considering the lower aging periods. Therefore, in order to compensate for this effect, the aging process are performed in liquid medium to accelerate the heating stage for the samples aged at 2 up to 20 min. The samples aged for times from 50 min up to 20, 000 min were aged in an air circulated oven. The termination of the carbon diffusion after aging was accomplished by quenching the samples in water after elapsing the aging-time. Three samples were investigated for each condition. The BH-response is defined as follows [24]:

8

1- For the non-pre-strained samples: Calculated from the difference between the lower yield (Rel) point after aging and the proof strength measured before aging (Rp0.2) as shown in Fig. 2a. 2- For the pre-strained samples: Calculated from the difference between Rel after prestraining and aging and the attained stress at the applied pre-strain (σx) as shown in Fig. 2b.

a

b

Fig. 2: Schematic tensile curves showing how the BH values are being estimated for samples (a) not pre-strain (BH0) and (b) with x% pre-strain (BHx).

It is worth mentioning that after heat treatments, after pre-straining and after aging the samples were directly kept in a freezer at a temperature of -20 °C until further processing. This is done to avoid natural ageing at room temperature. Moreover, for the same purpose, the samples after heat treatment were cleaned from the salt using cold water (not using warm water, which dissolves the salt much easier) then fatting to prevent oxidation before moving to the freezer. Figure 3 depicts the experimental sequence implemented during the current study.

9

Fig. 3: Schematic diagram depicting the experimental procedure adopted during this study. The annealing temperature T was varied to comprise six different Ti and a further annealing in the austenitic region.

3. Experimental Results and Discussions 3.1 The intercritical region In order to characterize the intercritical region, dilatometric investigation had been applied by heating the sample up to 950 °C. For assuring the equilibrium points, the dilatometric specimens were heated with a heating rate of 0.05 K/s. Using the lever rule, the transformed austenite fraction was calculated from the variation of the relative change in length as a function of temperature [17]. In Fig. 4 a comparison is given between the dilatometrically measured transformed austenite fraction with the predicted one obtained using the ThermoCalc software employing the TCFE8 database. The dilatometric curve starts to deviate from linearity at ~701 °C, indicating the beginning of the austenite formation. The curve-linearity is restored at ~855 °C, which corresponds to the end of austenite formation. Within this temperature-range (701 to 855 °C), the ferrite and austenite co-exist, i.e. this temperature-range corresponds to the intercritical region. With temperature-increase, the ferrite and austenite phases expand and simultaneously more 10

austenite is formed on the expense of the ferrite. The latter phase-transformation results in contraction of the sample because of the higher density of austenite compared to ferrite. At the beginning and end of the intercritical region the expansion due to temperature increase is dominant, whereas within the temperature range between 780 to 825 °C the contraction due to phase transformation dominates. The dominance of contraction within this temperature-range indicates the high rate of austenite formation within this range. Similarly, the calculated austenite fraction from the dilatometric curve applying the lever-rule shows that the transformation starts with a slow rate until it reaches a temperature of 780 °C and then continues rapidly till 825 °C. Nevertheless, it was expected from the ThermoCalc calculations that the austenite transformation begins at ~703 °C with a high rate till ~715 °C. During this temperature range the cementite-content decreases with increasing the temperature. Once the cementite phase disappears, the austenite formation continues at a much slower rate. With subsequent temperature-increase, the austenite formation rate increases. This deviation of the austenite formation rate, during dilatometric measurement, from the predicted one can be correlated to the sluggish kinetics of austenite formation during heating in the dilatometer close to Ae1, where Ae1 is the temperature separating the ferrite and ferrite + Austenite phase fields under equilibrium. Similar results had been reported by Meyer et. al [25]. A noticeable abrupt departure from the continuity of the dilatation curve (marked by the circle in Fig. 4) is observable within the intercritical region. This abrupt dilatation, that may be mistakenly correlated to carbide dissolution, does not indeed relate to kinetics of phase transformation but associates with the Curie temperature (TCu). Fig. 5 shows the distinct increase in the power required to keep the sample at the scheduled temperature, that appears when the sample transforms from the ferromagnetic to the paramagnetic state, i.e. at TCu is

11

correlated to the abrupt dilatation change. TCu of steel is known to be strongly dependent on the manganese concentration following the relation TCu = [769°C - XMn×15] °C,

(1)

where XMn is the Mn wt% [26]. This equation yields a value of 753.7 °C for TCu. The good consistence of the temperatures at which the heating power jumps, i.e. the abrupt dilation occurred, and the calculated TCu indicates that the temperature measurement system in dilatometer yields correct absolute temperature values.

Fig. 4 Predicted dependence of phases in equilibrium on the annealing temperature (using ThermoCalc) together with the formed austenite fraction calculated from the variation of the relative change in length in dilatometer.

12

Fig. 5 HF-power jump at Tcu together with the abrupt dilatation within the intercritical region.

3.2 The heat treatments The thermal cycles were monitored during heat treatments in the salt bath as well as during aging in the liquid medium using a data-logger. Representative thermal cycles are given in Fig. 6. It took about 25 s to approach the 170 °C, this interval is less than one quarter of time necessary to approach the same temperature when heating in air circulated oven. Fig. 6 illustrates that, 50 and 70 s are required to approach the annealing temperature of 760 and 900 °C, respectively. A cooling rate exceeding 1300 K/s in the temperature range 800 to 500 °C (temperature range, in which austenite can decompose to ferrite and/or bainite) was achieved. Quenching in brine from the annealing temperatures to room temperature was achieved within about 1 s. The achieved high cooling rate is due to using the brine, which yields a cooling rate as twice as that of water. Additionally, arranging the samples with keeping spacing among them (Fig. 1a) does not only assurance similar cooling rates in the samples but also results in a higher cooling rate. This high cooling rate guarantees freezing the microstructure obtained during annealing, i.e. suppressing the transformation of austenite into allotriomorphic ferrite or bainite during cooling until reaching the martensite start. Cooling of austenite phase from 13

the ferrite + austenite temperature range can substantially delay the transformation compared to cooling of the fully austenitized steel of the same composition. This is due to the enrichment of the intercritical austenite with C and Mn (as will be shown under section 3.3.4); accordingly, the bainite transformation is shifted to lower temperatures and slower cooling rates. Therefore, extremely fast cooling is necessary for high annealing temperatures, which corresponds to diluted alloying elements in austenite. In this context, Girina et al. observed no martensite, but a mixture of 25 % ferrite and 75 % bainite, in a 0.06C – 1.8Mn steel quenched from the fully austenitic region applying a cooling rate of 200 Ks-1. Nevertheless, a cooling rate of about 40 Ks-1 was enough for suppressing the bainite formation, when cooling the same steel from 760 °C (corresponds to 45 % austenite) and results in ferrite + martensite DP steel [27]. This study is accounted for investigating the dependence of the mechanical- and BH-response of DP steels on the MVF. This dependence would be affected if the martensite is partially replaced by bainite. Therefore, it was essential to suppress any bainite formation by quenching in brine. Although brine-quenching does not correspond to actual industrial condition, but it was necessary to fulfill the aim of this study. The extremely fast cooling applied in this study produces finer and more randomly oriented martensite structure than that industrially produced. The former structure would lead to elevated strength due to the resultant increased dislocation impedances [28]. Moreover, this high cooling rate illustrates an extreme bounding of carbon atoms in solid solution. More free carbon atoms increase their interaction with dislocations during the aging process and would result in higher BH-value [29]. On the other hand, in industrial process, high cooling rates are normally considered for the temperature range between 800 and 500 °C. The cooling at lower temperature is slower. Finally, the steel band is coiled, and the aging process starts in the coil.

14

Fig. 6 Monitored time-temperature regimes during heating in GS 540 salt-bath to annealing temperatures of 900 °C and 760 °C and quenching in brine and during heating in AS 140 salt-bath to aging temperature of 170 °C and quenching in water.

3.3 Microstructural features After the salt bath heat treatment, microstructure investigation was carried out. Light optical and scanning electron microscopes were used. 3.3.1 Recognition of martensite within ferrite phase Figure 7 reveals the valuable use of slight tempering for better identification of DP phases. Figures 7a and 7b correspond to samples without applying the slight tempering treatment. In Fig. 7a phases could not be fully identified under the LOM due to its small size and bad definition. SEM was used for a clearer identification of the phases. It is shown in Fig. 7b that the microstructure obviously appears clearer than that obtained from the LOM, but there is still uncertainty regarding some phases as those shown by arrows in this Figure. Consequently, a tempering treatment at 200 °C for 2 h was applied before etching. This slight tempering leads to a shallow darkening of the martensite in LOM as shown in Fig. 7c. Darkening of martensite by tempering followed by nital etching is a common technique adapted for simple microstructures to facilitate LOM examinations [30]. Decomposition products resulting from the slight tempering are the reason for darkening of the martensite

15

phase under the LOM, prolonged tempering at high temperature alters the martensite to spheroidized carbides in a matrix of ferrite [31]. The tempering procedure is also fitting for SEM observations [19]. This slight tempering leads to a clearer distinction of the phases. This can be seen when comparing Fig. 7b with 7d. Before tempering, martensite exhibits both a smooth and a featureless aspect. The tempering process furnishes the martensite islands with a structure because it induces a very fine carbide precipitation, which facilitate their ease distinguishing within the ferrite phase. Girault et al. [19] used this procedure to distinguishing martensite from retained austenite in the SEM micrograph. It was noticed also that after tempering carbides precipitates are to be perceived in the ferrite phase. Therefore, the slight tempered samples can be used for investigating the micro-characteristics of the DP. Of course, the nano-characteristics and the mechanical properties will be affected by such a treatment.

a

b

16

c

d

Fig. 7 (a) and (b) microstructures before tempering and (c) and (d) after tempering. The structures are for samples annealed at Ti = 760 °C.

3.3.2 Microstructure of the as-received DP600 Fig. 8 shows the as-received (AR) microstructure. LePera’s etchant was used to reveal the constituents in the optical microscopic investigation. The samples were tempered at 200 °C for 2 h before applying the etchant, this is to enable distinguishing the martensite from the retained austenite [19]. As shown in Fig. 8 a and b LePera’s etchant stained ferrite blue to light brown, martensite brown, and bainite grey. The retained austenite remains white. The AR-DPW600 contains (15.1 ± 2.9) % of the hard phase. The localization of bainite at one surface-side of the hot rolled sheet (Fig. 8b) indicates a slower cooling rate at this surface (lower side of the strip). Neither the middle (Fig. 8a and c) nor the other-surface reveal bainite structure. This may point out that the fast quenching is carried out from only one surface-side of the sheet (upper side). To confirm the occurrence of bainite at one side of the sheet, SEM investigation was performed. In this investigation no slight tempering is applied before nital etching. Excluding the slight tempering yields a featureless martensite after etching and hence ease distinguishing it from bainite in the SEM images. The bainite in Fig. 8d is characterized by its internal structure and it is surrounded by the featureless M/RA phase. The latter characteristic is also observable in Fig. 8b. 17

The mean grain size of the ferrite has a value of (3.45 ± 0.16) µm. The direction of hot rolling is being inherent in the microstructure, particularly in the developed martensite phase, which may indicate the formation of this microstructure from a pancaked, not fully recrystallized, parent austenite.

(b)

(a)

RA

M F

B B

F

M M RA

(c)

(d) F

M/RA M/RA

F

B

B

Fig. 8 LOM and SEM micrographs of the as received DPW 600 of (a) and (c) sample middle and (b) and (d) sample border. (a) and (b) LOM – slight tempered then etched with LePera. (c) and (d) SEM – non-tempered, etched with nital. F: Ferrite, M: Martensite, B: bainite and RA: Retained austenite.

3.3.3 Suppressing surface decarburization It was important for the current experiments to suppress the surface decarburization during heat treatment. Excessive decarburization is not tolerable, bearing in mind that the sample thickness is 2 mm and no surface machining will be carried out after the heat treatment. Therefore, surface-decarburization will result in a significant divergence in the results and leads to misleading conclusions. Therefore, performing the check-foil test directly before immersing the samples in the salt bath was crucial. The microstructures of check-foils 18

quenched in brine after immersing for 30 min at 800 °C in a neutral salt bath is shown in Fig. 9a and b and in a decarburizing one is shown in Fig. 9c. Fig. 9c demonstrates that decarburization of the foil resulted in formation of a ferrite phase instead of martensite. The formation of ferrite after quenching causes the high plasticity of the decarburized foil. On the other hand, the foil-decarburization is suppressed when immersing in the neutral salt bath. The structure of this foil is composed of martensite and carbide particles as shown in Fig. 9a and b. This brittle structure is the reason for fragmentation of the foil during bending between fingers, indicating the neutrality of the salt bath as described in section 2.2.3. In addition to the check-foil tests carried out during salt bath treatment, metallographic investigation of surface decarburization after each heat treatment were performed as well. A tolerable surface decarburization of up to 5 µm is observed as shown in Fig. 9c. Comparing the

sample-surface (Fig. 9c) with its middle (Fig. 9d) of a sample treated in a neutral salt bath

indicates the homogeneity of the microstructure.

a

b

19

c

d

Ti = 780 °C

e

Fig. 9 Microstructures of check-foils observed after immersing in a neutral salt bath (a) and in a decarburizing one (b). Surface (c) and middle (d) of a sample annealed at Ti = 800 °C in a neutral salt bath.

3.3.4 Developed phases In Fig. 10, LOM and SEM photomicrographs are shown for each annealing temperature applied in this investigation. The different annealing temperatures applied during the heattreatment processes to produce DP steels with different MVFs, are based on the results given in Fig. 4. It is clearly observable in Fig. 10 that increasing the annealing temperature resulted in increasing the MVF and decreasing the ferrite grain size (df). Table 2 lists both of the MVF measured applying the point count method and the df measured applying the mean linear intercept method. Table 2: Dependence of the MVF and df on the annealing temperature T [°C] MVF [-] df [µm]

740 0.169 4.44

760 0.224 3.72

780 0.300 3.99

800 0.388 3.45

825 0.545 3.36

840 0.864 2.63

20

It is obvious that the martensite type formed when annealing at 740, 760, 780 and 800 °C is mainly the plate-type martensite. Annealing at 825, 840 °C and 900 °C results in the formation of a mixture of the plate- and the lath-type martensite. This change in martensite type is correlated to the variation in the austenite-composition from which the martensite is formed as shown in Fig. 11. The austenite transformed basically to lath martensite, when it forms from austenite with diluted alloying elements (like that formed at 825, 840 and 900 °C). Each lath is a result of a homogeneous shear, and successive shears produce a packet of parallel laths containing a high density of tangled dislocations. Basic units are generally aligned parallel to one another in packets [33]. The martensite packets are made up of units ranging from less than 0.1µm to several microns. Plate martensite differs from lath martensite in that its adjacent plates do not form parallel to one another. The plates which are the first to form tend to span their parent austenite grains. Smallman et. al. reported that the martensite structure found in dual-phase steels has the

characteristics of plate martensite [34]. The observation of the lath martensite as the main structure of the DP steel annealed at 840 °C is correlated to the fact that a DP steel with MVF of 0.86 is not a common one. Referring to Fig 11, it is clear that heating into the two-phase region substantially increases the austenite stability via enrichment of austenite with the hardenability-increasing elements, namely the C, Mn and Cr. This is an extremely important aspect of the DP steel, which allows for the intercritical austenite to undergo transformations that are mostly typical for medium carbon steels instead of the low carbon content of its alloy. In this transformation the austenite is so stable that it bypasses the ferrite/pearlite and the bainite noses of the continuous cooling transformation diagram and therefore the austenite undergoes martensite transformation. Nevertheless, with increasing the austenite-content, i.e. increasing the annealing temperature, the austenite stabilizing elements decrease as shown in Fig. 11. During the current study, suppressing the ferrite/bainite transformation during quenching was possible, even in DP 21

steels annealed at high temperature, thanks to the applied extremely high cooling rate as shown under section 3.2.

MVF = 0.169

Ti = 740 °C Ti = 740 °C MVF = 0.224

Ti = 760 °C Ti = 760 °C

22

MVF = 0.300

Ti = 780 °C

Ti = 780 °C

MVF = 0.388

Ti = 800 °C Ti = 800 °C MVF = 0.545

Ti = 825 °C Ti = 825 °C

23

MVF = 0.864

Ti = 840 °C Ti = 840 °C MVF = 1.0

900 °C 900 °C Fig. 10 LOM (left) and SEM (right) micrographs obtained after annealing at the prescribed annealing temperature.

24

Fig. 11 Dependence of the austenite composition on the austenite volume fraction (transforms via quenching to martensite) – predicted from ThermoCalc calculations.

Samples annealed in salt baths at 740, 760, 780, 800, 825 and 840 °C and quenched in brine down to room temperature were taken to measure the volume fraction of the martensite using the systematic point count method. For a better distinguishing of the martensite phase, the volume fraction counting method was performed on images of SEM. The results are presented in Fig. 12 together with the predicted (Thermo-calc) and calculated (dilatometric) austenite volume fraction. In this figure, the comparison of the volume fraction of martensite with the austenite volume fraction is based on the assumption that the entire intercritical austenite transforms to martensite during quenching from the intercritical region. This assumption is a realistic one because of the very high cooling rate applied during this study. It is observed that the measured MVF gradually deviates from the predicted one with increasing the annealing temperature. The measured MVFs for Ti of 740, 760 and 780 °C have similar values as the predicted ones using the ThermoCalc software, whereas the measured MVFs for Ti of 800, 825 °C and 840 °C are significantly lower than the predicted ones. The MVFs calculated 25

applying the lever-rule on the dilatometric curve show similar values to the measured one only for Ti of 800°C. The lever rule does not give an accurate transformed phase amount because it does not take into account that the carbon redistribution between the forming ferrite and the remaining austenite. Therefore, it cannot distinguish between the dilatation experienced by the specimen due to the formation of ferrite and the dilatation experienced by the austenite due to progressive carbon enrichment [25, 35].

Fig. 12 Comparison of dependence of austenite volume fraction predicted by ThermoCalc and that calculated from the dilatometric measurement with the MVF measured in samples annealed in salt bath before quenched. Please note that the comparison between the austenite volume fraction and the MVF is based on the assumption that the entire intercritical austenite is transformed to martensite during quenching to room temperature.

3.4 Mechanical behavior Figure 13 shows the engineering stress–strain curves for the different treatment conditions. Increasing the intercritical MVF results in increasing the strength of the DP steel on the expense of its ductility. Figure 14 summarizes this effect together with the presentation of the properties obtained in the as-received DPW600. The adopted heat treatment approach resulted in a high reproducibility of the tensile properties as can be inferred from the small error bars 26

of the diagram. However, the reproducibility of the total elongations (TEl) are not as high as that of the other tensile properties (TEl shows larger error bars). This results divergency is attributed to the geometric instability after necking during tensile testing and not to the reproducibility of the heat treatment cycle. It is shown in this diagram that dependences of the ultimate tensile strength (Rm) and the 0.2% proof strength (Rp0.2) on MVF can be fitted by a linear function. Based on the law of mixtures one can write the Rm and R p0.2 of DP steel as a function of MVF as follow: Rm = RmM × MVF + RmF × (1-MVF) = RmF + (RmM - RmF) × MVF,

(2)

Rp0.2 = Rp0.2M × MVF + Rp0.2F × (1-MVF) = Rp0.2F + (Rp0.2M – Rp0.2F) × MVF,

(3)

where RmF and RmM are the ultimate tensile strength of ferrite- and martensite phase of the DP steel, respectively. Rp0.2F and Rp0.2M are the yield strength of ferrite- and martensite phase of the DP steel, respectively. In case that RmF, RmM, Rp0.2F and Rp0.2M are invariant with respect to the MVF, the tensile strength of a DP steel can be written as Rm = Am + Bm × MVF,

(4)

Rp0.2 = Ap0.2+ Bp0.2 × MVF,

(5)

where Am, Bm, Ap0.2 and Bp0.2 are constants. Therefore, the observed linear dependence of the strength on MVF indicates that the rate of strengthening of DP steel is independent on the intrinsic strength of its ferrite and martensite phases. Indeed, the martensite strength is predicted to decrease with increasing MVF because its increase results in concentration reduction of the martensite strengthening elements as shown in Fig. 11. However, Davis concluded that the strength of a dual phase structure is independent of the composition and strength of the martensite [7]. 27

The here observed linear relationship between the strength of DP streel and its MVF is also reported by other researchers [7, 9, 10, 11, 12]. Byun and Kim [9] reported Am and Bm values of equation (4) to be 677 and 564, respectively. These values are in a good consistence with the observed factors of this study (see Fig. 14). However, results presented in others work show no linear relation between the strength of DP and MVF [8, 36, 37, 38]. Birgani and Pouranvari reported that Rp0.2 linearly increases by increasing MVF, while Rm increases with increasing MVF, peaking around 50 % and decreasing with further increase in MVF [36]. Similar MVF peaking value is reported for both Rm and Rp0.2 by Bags and Rays [38]. However, they showed an increase in strength until 100 % martensite following the last decrease after peaking. According to equation (2) and (3), the variation of RmF, RmM, Rp0.2F and Rp0.2M with MVF can cause the nonlinear dependence of the Rm and Rp0.2 on MVF. Bag and Ray [37] expressed this variation as a polynomial function of the inverse of the mean free path, λ-1, which in turn varies by varying the MVF. They considered the mean free path as the most important factor affecting RmF and RmM because the development and distribution of all types of stresses in ferrite or martensite in DP steel depend on this parameter. Correspondingly, a model justifying the unusual behavior of strength variation with varying MVF has been developed [37]. Similarly, Chang and Preban [39] found that the strength of their studied DP steels varied linearly with inverse square root of mean free path in ferrite. The reason for the contradicting reports regarding the dependence of RmF and RmM on MVF in [8, 36, 37, 38, 39] on the one hand, and their independence on MVF in the current study and in [7, 9, 10, 11, 12] on the other hand is vague and needs further investigation.

28

Fig. 13 Engineering stress-strain curves at the prescribed MVFs

Fig. 14 Effect of MVF on the tensile properties together with the properties of the as-received (AR) DPW600

On the other hand, the uniform elongation (UEl) and the total elongation (TEl) progressively decrease with increasing the martensite. The dependence of (TEl) and the uniform elongation (UEl) on MVF were fitted with a second-degree function as shown in Fig. 14.

29

High ductility at given strength level is one of the greatest advantages of DP steels. The uniform elongation at high strength plays an important role in the sheet metal forming. The UEl relates well with stretch-forming operations. In these operations, the material is considered to fail after necking, which corresponds to the end of the uniform elongation. However, the TEl is a decisive factor, when considering other manufacturing operations, such as bending and countersink. The non-linear decrease of UEl and TEl with increasing MVF indicates the presence of another playing-factor, which counters the effect of increasing the MVF on decreasing the ductility. This factor could be the martensite softening with increasing MVF (i.e. increasing Ti) due to the continual concentration-decrease of alloying elements as previously shown in Fig. 11. The higher incompatibility of phases-deformation at higher relative strength of martensite (i.e. lower MVF) results in a premature formation of voids in steels and hence lower ductility [40]. This incompatibility of phases-deformation between martensite and ferrite reduces by increasing the martensite content and consequently softening the martensite. The effect of increasing the MVF on ductility dominates, therefore the TEl decreases until shortly before reaching the 100% martensite. The AR DPW600, which contains MVF of 0.151, showed significantly higher Rp0.2 than the values calculated by the fitting functions (Fig. 13). This behavior is typical for the aging effect. The tensile samples used in this investigation are produced out of shelved sheets stored for about 15 years at room temperature. The natural aging effect taking place during this long storing time resulted in changing the mechanical properties. Substituting the MVF of the ARDPW600 in the Rp0.2-equation given in Fig. 14 yields a predicted value of Rp0.2 of 266 MPa, which is about 88 MPa lower than the measured one.

30

3.5 Bake hardening behavior The dependence of the BH-response on the aging time is given in Fig. 15. The changes in the tensile stress-strain-relationship during ageing develop first by locking the dislocations due to the formation of the Cottrell effect [41]. The first increase in the yield strength in the samples aged at 100 and 170 °C (Fig. 15 c-f) is believed to be due to the development of this atmosphere. The samples treated at 220 °C (Fig. 15 a and b) did not show this stage because of the fast aging-kinetics at this temperature, such that aging-time of two minutes was enough for the development of the second aging stage. The second strengthening-stage, which ultimately results in an appearance of upper plateau in the BH-response curves is due to the precipitation of carbides at dislocations. Both stages lead to the pinning of dislocations and hence strengthening of the steel. An impressive increase in the yield strength due to aging during the second stage is observable. However, for the DP aged at 100 °C, the development of the upper plateau is only observable in the samples treated at 840 °C as shown in Fig. 15 f. This is attributed to the sluggish diffusion rate at 100 °C. The high dislocation density in ferrite and the lower diffusion distance expected in the DP annealed at 840 °C with MVF = 0.86 is the reason of the fast developed second plateau, instead of the low aging temperature. At longer ageing times, an over-aging process can be observed, since the yield strength drops sharply once the maximum has been reached. This behavior was observable only on the samples aged at 170 °C and 220 °C (Fig. 15 a-e). 3.5.1 Effect of pre-strain The BH response, due to the formation of Cottrell atmospheres, is strongly dependent on prestraining. The pre-straining has not been applied on the samples with Ti = 840 °C (MVF = 0.86) because these samples will approach the necking point at 2 % pre-strain and will be past it at 5 % as shown in Fig. 13. 31

Fig. 15 c and f demonstrates that the first strengthening-stage, due to formation of Cottrell atmosphere, exhibits values ranging from ~20 to ~40 MPa, when considering the non-prestrained samples. For the 2 % and 5 % pre-strained ones, the Cottrell atmosphere formation results in an increase in the BH-response of ~55 to ~85 MPa as presented in Fig. 15 d-f. The progress in the development of Cottrell atmosphere is obvious only for the investigated nonpre-strained samples annealed at Ti = 760 °C and aged at 100 °C as shown in Fig. 15f. For the other tested aging-temperatures, the minimum studied aging time of 2 min was enough for development of Cottrell atmosphere. Further segregation of solute has little or no effect on the dislocation pinning, this results in the appearance of the plateau in the diagram after developing the Cottrell-atmosphere. An outright plateau is not observable in many aging conditions but it is obscured by fluctuation around a certain value. During the subsequent stage of ageing, continuing segregation causes occurrence of precipitates formed in the later stages of ageing. In contrast to the effect of pre-straining on BH-response during the first strengthening stage, it is obvious during the second stage that pre-straining has a negative effect on it. Therefore, the maximum achieved BH-response decreases with increasing the pre-straining value. 3.5.2 Effect of aging temperature It is evident that the aging temperature has a significant effect in accelerating the aging process. The different stages of aging occur after shorter aging times with increased aging temperatures. For aging temperature of 220 °C, aging time of 2 min was enough for the development of the second aging stage (Fig. 15 a and b). This stage developed within aging time between 50 and 500 min (depending on pre-strain value and MVF) when aging at 170 °C as shown in Fig. 15 c-e. Decreasing the aging temperature to 100 °C results in further deceleration of the aging kinetics that the development of the second stage requires more than 10 000 min. for Ti = 760 °C and about 1000 min. for Ti = 840 °C as shown in Fig. 15 f. It is notable here that the BH level of the first and the second stages are independent of the aging 32

temperature, and they are only determined by the MVF (i.e. Ti) and the pre-strain level. Accordingly, changing the aging temperature does not change the characteristics of the agingphenomenon but it affects its rate.

3.5.3 Effect of MVF It can be inferred from Fig. 15 b, c and e that the DP annealed at 740 °C, which has the lowest MVF of 17 % recorded inferior BH-response during both of the first and the second strengthening stage. The samples annealed at 760, 780, 800 show similar level of BHresponse, when considering constant BH-parameters. Fig. 15 f shows that the samples annealed at 840 °C with only about 14 % ferrite reveal similar level of strengthening due to the formation of Cottrell atmosphere like the corresponding samples annealed at 760 °C. Nevertheless, the carbide-precipitation process proceeds in the former samples much faster instead of the low aging temperature of 100 °C to a value of about 85 MPa. The latter value is approximately 1/2 of the BH level achieved in the non-pre-strained DP steels annealed at 760 °C and 800 °C (Fig. 15 a and c). Please note that the latter comparison is based on the observation reported in the previous section that the aging temperature has an insignificant effect on the BH level of the first and the second plateaus. Accordingly, the strengthening stage due to carbide precipitation increases with increasing MVF, reaching a saturation level of BH-response at a certain range of MVF before decreasing. In the current study this saturation range is observed to be between 0.22 and 0.39.

33

a

b

c

d

e

f

Fig. 15 Dependence of the BH-response on the aging time of the DP steels for: (a) and (b) aging temperature of 220 °C, (c), (d) and (e) aging temperature of 170 °C and (f) aging temperature of 100 °C. The investigated Ti and prestraining value for each aging temperature are prescribed in the diagram.

4. Conclusions Tensile properties and strain aging behavior of ferrite-martensite dual phase steels containing 0.17-1.0 martensite volume fraction was analyzed. From the results of the study, the following conclusions are obtained: 1- Prediction of the MVF using thermodynamic-calculations with ThermoCalc were appropriate up to a value of 0.3, which corresponds to intercritical annealing at 780 34

°C. The obtained MVFs at higher annealing temperatures were significantly lower than the predicted ones. On the other hand, the transformed phase-amounts calculated applying the lever rule on the dilatometric curve deviates significantly from both the predicted and the actually measured MVFs. 2- The Rm and Rp0.2 values show a linear dependence on MVF. This linearity indicates that the strength of the studied DP steels is independent of the intrinsic strength of martensite, which is predicted to increase with decreasing MVF. The latter predicted strength-increase is attributed to the increase in alloying elements concentration in the intercritical austenite, from which the martensite is formed. 3- With increasing the aging time, the DP steel undergoes two strengthening-stages due to the BH-effect. These are an initial increase in yield strength in the first stage (due to Cottrell atmosphere formation) followed by a larger yield strength increase in the second one (due to carbide-precipitation). Prolonging the aging, after attaining the second stage, results in decreasing the BH-response achieved in the second stage. 4- Increasing the aging temperature does not result in changing the strengtheningcharacteristics described in the previous conclusion but it accelerates its rate. Thus, for similar MVF and pre-strain value, the BH-responses achieved in the first and second strengthening-stages are insignificantly affected by the aging temperature. 5- By contrast to conclusion number four, pre-straining is a decisive factor for the BHresponse. Pre-straining the DP steel significantly increases the BH-response of the first strengthening stage, but it negatively influences it in the second one. 6- In the current study it is shown that, under constant strain-aging parameters, varying the MVF between 0.22 and 0.39 has insignificant effect on the strengthening due to BH-effect. Decreasing MVF to 0.17 and increasing it to 0.86 results in decreasing the BH-response of the second strengthening stage.

35

Acknowledgements The authors would like to thank ThyssenKrupp Steel Europe AG for supplying the steel as well as the Deutsche Forschungsgemeinschaft (DFG), Germany, financially supporting the investigation [grant number PA 837/33-1].

Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

References [1]

M. Calcagnotto, D. Ponge, E. Demir, D, Raabe, Orientation gradients and geometrically necessary dislocations in ultrafine grained dual-phase steels studied by 2D and 3D EBSD, Mater. Sci. Eng., A 527 (2010), 2738–2746.

[2]

A.H.

Nakagawa,

G.Thomas,

Microstructure-mechanical

property

relationships of dual-phase steel wire, Metall. Trans. A 16 (1985), 831– 840. [3]

D. Ji, M. Zhang, D. Zhu, S. Luo, L. Li, Influence of microstructure and pre-straining on the bake hardening response for ferrite-martensite dualphase steels of different grades, Mater. Sci. Eng., A 708 (2017), 129–141.

[4]

H.K.D.H. Bhadeshia, D.V. Edmonds, Analysis of mechanical properties and microstructure of high-silicon dual-phase steel, Metall. Sci. 14 (1980) 41–49.

[5]

N. Fonstein, H-J Jun, G. Huang, S. Sadagopan, B.Yan, Effect of bainite on mechanical properties of multi-phase ferrite-bainite-martensite steels, In MS&T’11, Columbus, Ohio, (2011) 634-641.

[6]

R.D. Lawson, D.K. Matlock, and G. Krauss, Fundamentals of Dual-Phase Steels, R.A. Kot and B.L. Bramfitt, eds., TMS-AIME, Warrendale, PA, (1981) 347–381.

[7]

R.G. Davies, Influence of martensite composition and content on the properties of dual phase steels, Metall. Trans. 9A (5) (1978) 671–679.

[8]

B. Aksakal, F. Karaca and R. Arikan, Effect of martensite volume fraction on mechanical properties of dual-phase treated AISI-4012 sheet steels, 36

Int. J. Mat. Res. (formerly Z. Metallkd.) 101 (2010), 684–691. [9]

T.S. Byun, Kim I.S., Tensile properties and inhomogeneous deformation of ferrite-martensite dual-phase steels, J. Mater. Sci., 28 (1993), 2923– 2932.

[10] R.G. Davis, C.L. Magee, The deformation behavior of a vanadiumstrengthened dual phase steel, Metall. Trans. 9A (1978) 41–52. [11] A.N. Bortsov, N.M. Fonshteyn, The distribution of strains between the phases of ferrite martensite steels, Phys. Met. Metall., 61 (1987), 74–81. [12] N. Fonstein, Dual-phase steels, in:

R. Rana, S. B. Singh (Eds),

Automotive Steels, Design, Metallurgy, Processing and Applications, Woodhead Publishing - Elsevier Ltd, 2017, 169–216. [13] T. Anke, Doctoral dissertation: Bake Hardening von warmgewalzten Mehrphasenstählen, Pieper, Clausthal-Zellerfeld, 2005 [14] A.J. Abdalla, L.R.O. Hein, M. S. Pereira, T.M. Hashimoto, Mechanical behaviour of strain aged dual phase steels, Materials Science and Technology 15 (1999), 1167–1170. [15] T. Waterschoot, B.C. de Cooman, A.K. De, S. Vandeputte, Static strain aging phenomena in cold-rolled dual-phase steels, Metall and Mat Trans A 34 (2003), 781–791. [16] M. Soliman, B. Weidenfeller and H. Palkowski, Metallurgical Phenomena during Processing of Cold Rolled TRIP Steel, Steel Res. Int. 80 (2009), 57-65. [17] M. Soliman and H. Palkowski, On Factors Affecting the Phase Transformation

and

Mechanical

Properties

of

Cold-Rolled

Transformation-Induced-Plasticity–Aided Steel, Metall and Mat Trans A, 39 (2008), 2513–2527. [18] Victoria Cain, High temperature creep behaviour niobium bearing ferritic stainless steels, Thesis, Cape Peninsula University of Technology (2005), P. 62 [19] E. Girault, P. Jacques, Ph. Harlet, K. Mols, J. Van Humbeeck, E. Aernoudt, F. Delannay: Metallographic Methods for Revealing the Multiphase Microstructure of TRIP-Assisted Steels, Mater. Char. 40 (1998), 111–118. 37

[20] ASTM E562 – 19, Standard Test Method for Determining Volume Fraction by Systematic Manual Point Count, ASTM Int., PA, 2019. [21] ASTM E112 – 13, Standard Test Methods for Determining Average Grain Size, ASTM Int., PA, 2013. [22] ASTM E8-13a, Standard test methods for tension testing of metallic materials, ASTM Int., PA, 2013. [23] Norm DIN EN ISO 6892-1: Metallische Werkstoffe – Zugversuch – Teil 1: Prüfverfahren bei Raumtemperatur, Beuth Verlag, Berlin, 2009. [24] DIN EN 10325, Stahl – Bestimmung der Streckgrenzenerhöhung durch Wärmebehandlung (Bake-Hardening-Index), Deutsche Fassung EN 10325, Beuth Verlag, Berlin, 2006. [25] M. De Meyer, J. Mahieu, B.C. De Cooman, Empirical Microstructure Prediction Method for Combined Intercritical Annealing and Bainitic Transformation of TRIP steel, Mater. Sci. Technol., 10 (2002), 1121– 1132. [26] T.A. Kop, J. Sietsma, S. Van der Zwaag: Dilatometric Analysis of Phase Transformations in Hypo-eutectoid Steels, J. Mater. Sci., 36 (2001), 519– 526. [27] O.A. Girina, N.M. Fonstein, D. Bhattacharya: Effect of annealing parameters on austenite decomposition in a continuously annealed Dualphase steel, MS&T'03, Chicago, Illinois, (2003), 45th MWSP, Vol. XLI, 403–414. [28] A. Bardelcik, C.P. Salisbury, S. Winkler, M.A. Wells, M.J. Worswick, Effect of cooling rate on the high strain rate properties of boron steel of boron steel, Int. J. Impact Eng., 37 (2010), 694–702. [29] A. Momeni, K. Dehghani, S. Abbasi, M. Torkan, Bake hardening of a low carbon steel for automotive applications, Metalurgija 13 (2007), 131–138. [30] G.R. Speich, V.A. Demarest, and R.L. Miller, Formation of austenite during intercritical annealing of dual-phase steels, Metall. Trans. A (1981), 1419 – 1428. [31] Michael Bauccio - ASM Metals Reference Book, 3rd Edition, ASM International, Materials Park, OH, 1993. 38

[32] M. Soliman, H. Palkowski, Influence of hot working parameters on microstructure evolution, tensile behavior and strain aging potential of bainitic pipeline steel, Mater. Des. 88 (2015) 759–773. [33] G. Krauss, A.R. Marder, The morphology of martensite in iron alloys, Met.Trans. 2 (1971), 2343–2357. [34] R.E. Smallman, A.H.W. Ngan, in Modern Physical Metallurgy (Eighth Edition), Amsterdam, Butterworth-Heinemann, 2014 [35] M. Gómez, S.F. Medina, G. Caruana, Modelling of phase transformation kinetics by correction of dilatometry results for a ferritic Nb microalloyed steel, ISIJ Int., 43 (2003), 1228–1237. [36] E. N. Birgani, M. Pouranvari, Effect of martensite volume fraction on the work hardening behavior of dual phase steels, Metal 2009 (2009) Hradec nad Moravici. [37] A. Bag and K.K. Ray, A new model to explain the unusual tensile behavior of high martensite dual-phase steels, Metall and Mat Trans A, 32 (2001), 2400–2403. [38] A. Bag, K.K. Ray, and E.S. Dwarakadasa, Influence of martensite content and morphology on tensile and impact properties of high-martensite dualphase steels, Metall and Mat Trans A, 30 (1999), 1193-1202. [39] Peng-Heng Chang and A. G. Preban, The effect of ferrite grain size and martensite volume fraction on the tensile properties of dual phase steel, Acta Metall. Vol. 33, no. 5, pp. 897-903, 1985. [40] G.R. Speich, R.L. Miller, Mechanical properties of ferrite-martensite steels, in: R.A. Kot and J.W. Morris (Eds.) Conference Proceedings Structure and Properties of Dual-Phase Steels, New Orleans, LA, 1979. [41] A.H. Cottrell, B.A. Bilby, Dislocation theory of yielding and strain ageing of iron, Proc. Phys. Soc. A (1949) 49–62.

39

CRediT author statement Mohamed Soliman: Conceptualization, Methodology, Software, Validation, Formal analysis, Investigation, Resources, Data Curation, Writing - Original Draft, Visualization, Supervision, Project administration, Funding acquisition. Heinz Palkowski: Conceptualization, Resources, Writing - Review & Editing, Supervision, Funding acquisition.

Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: