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MECHANICAL PROPERTIES, SELECTIVE OXIDATION AND REACTIVE WETTING
OF A MEDIUM-MN THIRD GENERATION ADVANCED HIGH STRENGTH STEEL
K.M.H. Bhadhon1, T. Sydor1, J.R. McDermid1*, F.E. Goodwin2, A.P. Domingos3
1 Centre for Automotive Materials and Corrosion, McMaster University, Hamilton, ON, L8P 0A6, Canada
2 International Zinc Association, Durham, North Carolina, 27713, United States of America
3 International Zinc Association, Brussels, Belgium
* mcdermid@mcmaster.ca
ABSTRACT
A prototype medium-Mn third generation advanced high strength steel (3G AHSS, 0.2C-6.25Mn-
1Al-1Si-0.5Cr (wt.%)) with a martensitic starting microstructure was subjected to a range of
continuous galvanizing line (CGL) compatible intercritical annealing (IA) heat treatments. The
mechanical property evolution of the heat-treated steels showed that gradual transformation of the
retained austenite in the 680 °C × 120 s sample produced the highest UTS × TE product (29,800
MPa%), meeting 3G AHSS target mechanical properties. These IA parameters were used for
selective oxidation and reactive wetting studies. A multi-micron deep solute-depleted layer –
produced during the austenitizing heat treatment – and an intermediate flash pickling step,
conducted between the austenitizing and IA – resulted in a pre-immersion surface comprising
extruded Fe nodules, large area fractions of boldly exposed substrate and widely spaced nano-scale
oxide nodules. During continuous galvanizing in a conventional 0.2 wt.% Al (dissolved)
galvanizing bath, direct reactive wetting of the exposed substrate and Fe nodules along with
multiple secondary reactive wetting mechanisms resulted in a high-quality, robust galvanized
coating. Overall, the CGL-compatible thermal processing route used in this study was successful in
producing 3G AHSS target mechanical properties as well as robust galvanized coatings for the
prototype medium-Mn steel.
KEYWORDS
Medium-Mn 3G AHSS, mechanical properties, selective oxidation, reactive wetting, continuous
galvanizing
1. INTRODUCTION
Medium-Mn steels have great promise amongst the third-generation advanced high strength
steels (3G AHSS) to meet vehicle weight reduction targets whilst maintaining vehicle safety
through their improved target mechanical properties (i.e., 24,000 MPa% ≤ UTS × TE ≤ 40,000
MPa%) [1] versus those of the current first generation AHSS. However, a major challenge for the
successful integration of 3G AHSS into the automotive industry involves developing continuous
galvanizing line (CGL) compatible thermal processing routes. This implies the target mechanical
properties need to be achieved through heat treatment in the atmosphere-controlled furnace section
of the CGL while simultaneously producing a robust, high-quality galvanized coating.
Recently, numerous researchers working on medium-Mn steels have been successful in
achieving 3G AHSS target mechanical properties [2– 6]. However, these studies tended to use long
soaking times (≥ 120 – 180 s), making them incompatible with CGL processing parameters.
Furthermore, these studies were confined to mechanical property development and did not
investigate the selective oxidation and reactive wetting studies required to evaluate their
galvanizibility. Research on the various generations of AHSSs has shown that the commonly used
13th International Conference on Zinc & Zinc Alloy Coated Steel Sheet (GALVATECH 2023)
October 15~19, 2023 / COEX, Seoul, Korea
122
AHSS alloying elements – e.g., Mn, Si, Al, and Cr – selectively oxidize in the dew point controlled
N2-(5-20)H2 (vol%) process atmospheres used in the CGL. Depending on the oxide chemistry,
morphology, and distribution, these external oxides can lead to poor reactive wetting and
unacceptable coating quality during continuous galvanizing [ 7– 10]. Hence, it is important to
determine the effect of process atmosphere oxygen partial pressure (pO2) during annealing on the
selective oxidation and reactive wetting of medium-Mn steels.
In this regard, the present authors have been successful in achieving 3G target mechanical
properties with three prototype medium-Mn steel compositions in the Fe-(0.15-0.20)C-6Mn-xSi-
yAl-(0-0.05)Cr (wt.%) system using CGL-compatible thermal processing cycles [11– 13]. The
authors reported that a martensitic starting microstructure was more robust in achieving the target
mechanical properties using annealing times compatible with the CGL owing to the more rapid
formation of chemically and mechanically stable retained austenite during the intercritical annealing
(IA) heat treatment versus using the as-received tempered martensite microstructure [11–14]. This
retained austenite transformed to martensite gradually during deformation and aided in sustaining
high work hardening rates (WHR) by continuously supplying fresh barriers (i.e., geometrically
necessary dislocations and fresh martensite), thereby reducing the dislocation mean free path.
Furthermore, deformation induced nano-twin formation in the retained austenite was also observed,
suggesting twinning induced plasticity (TWIP) was working alongside the transformation induced
plasticity (TRIP) mechanism to provide a high strength/ductility balance in the prototype medium-
Mn steels [11–13]. Selective oxidation [15,16] and reactive wetting [17] studies were also carried
out for promising heat treatment cycles to determine the galvanizibility of the prototype medium-
Mn steels. In summary, high-quality galvanized coatings were obtained through successful reactive
wetting of the two medium-Mn steels investigated in [11,13] using CGL-compatible process
atmosphere and galvanizing parameters [17].
As a continuation of developing a CGL-compatible thermal processing route for a wide range of
medium-Mn steel chemistries, this study focuses on the development of 3G AHSS mechanical
properties, selective oxidation, and reactive wetting of a 0.2C-6.25Mn-1Al-1Si-0.5Cr (wt.%)
medium-Mn steel, building on the work documented in [11–17]. The main objective of this research
is to determine the robustness of CGL-compatible thermal processing parameters in producing 3G
AHSS target mechanical properties along with a high-quality galvanized coatings for the prototype
medium-Mn steel.
2. MATERIAL AND EXPERIMENTAL METHODS
The chemical composition of the prototype medium-Mn (med-Mn) steel is listed in Table 1. The
steel was cast as ingots in a vacuum induction furnace at U.S. Steel R&D (Munhall, PA). The ingots
were reheated to 1250 °C and hot rolled (HR) to a thickness of 25 mm with a HR finish temperature
of 1050 °C. The roughed ingots were slow cooled in vermiculite. The roughed ingots then
underwent a second HR step in which they were reheated to 1250 °C and rolled to a thickness of
4 mm with a HR finish temperature of 900 °C. The hot rolled steels were coiled at 665 °C and
isothermally held for 2 h, followed by slow cooling (7 °C/h) to room temperature. Subsequently,
18 mm was trimmed from each edge, followed by surface grinding to an ingot thickness of 2.7 mm.
Finally, the ingots were cold rolled to a final thickness of 1.2 mm.
Table 1. Chemical composition of the prototype medium-Mn 3G AHSS (wt.%)
C
Mn
Al
Si
Cr
Mo
Si/Mn
0.20
6.25
1.08
1.04
0.51
0.031
0.17
123
Two sample geometries – 10 mm × 50 mm coupons and 120 mm × 200 mm panels, where the
longitudinal axis was parallel to the rolling direction (RD) – were used in this study. The 10 mm ×
50 mm coupons were used for the selective oxidation study whereas the 120 mm × 200 mm panels
were used for the mechanical property development and galvanizing studies. The selective
oxidation coupons were ground with a series of SiC papers, with 1200 grit SiC being the final step,
to remove the effect of surface roughness on the subsequent surface analyses. All heat treatments
were carried out in the McMaster Galvanizing Simulator (Iwatani-Surtec) under a N2–5 vol.% H2
process atmosphere with a controlled dewpoint – i.e., controlled pO2. The temperature of the steel
samples was monitored by a type K thermocouple (0.5 mm) spot welded to the sample prior to the
heat treatments. The cooling rate of the annealed samples was controlled by N2 gas flow which was
adjusted to meet the targeted cooling rates via feedback control.
Fig. 1. Schematic diagram of the thermal processing cycle.
Table 2. Austenitizing and IA heat treatment parameters.
Heat Treatment Temp.
(°C)
Holding
Time (s)
OT
(°C)
Holding
Time (s)
Tdp
(°C)
pO2
(atm)
Austenitizing
(M)
890 600 –10 1.03 × 10-19
Intercritical
Annealing
(IA)
680
120 460 30
–30
3.96 × 10-26
–10
1.26 × 10-24
700
–30
1.43 × 10-25
720
–30
4.92 × 10-25
Fig. 1 shows a schematic diagram of the two-stage thermal processing cycle used in this study
[11–13]. In all cases, an austenitizing heat treatment was carried out at 890 °C for 600 s followed by
gas quenching at –30 °C/s to room temperature to produce the desired martensitic (M) starting
microstructure. A process atmosphere dew point of –10 °C was used during the austenitizing heat
treatment. For the mechanical property development study, a range of intercritical annealing (IA)
temperatures were chosen based on dilatometry data. Table 2 lists all heat treatment parameters
124
used. For the selective oxidation and galvanizing trials, all samples were flash pickled prior to the
IA treatment to remove the thick external oxide formed during the austenitizing. Flash pickling
consisted of a 60 s immersion in a 30 °C solution comprising 64.5 mL DI water, 59.6 mL HCl and
0.25 g hexamethylenetetramine. The flash pickled M samples were intercritically annealed at
680 °C for 120 s – the IA parameters that resulted in the highest UTS × TE product – followed by
cooling at –10 °C/s to an overaging temperature (OT) of 460 °C, where the samples were held for
30 s. The overaging section of the heat treatment was used to establish thermal equilibrium in the
steel substrate prior to entering the 460 ℃ molten Zn bath. Finally, the samples were cooled to
room temperature at –20 °C/s. For galvanizing trials, the steel panels were immersed in a Fe-
saturated 0.2 wt.% Al (dissolved) galvanizing bath for 4 s at 460 °C after the OT treatment. Two
process atmosphere dew points (Tdp) – –30 °C and –10 °C – were used during IA treatments to
determine the effect of process atmosphere pO2 on the selective oxidation and subsequent reactive
wetting of the prototype medium-Mn steel.
Room temperature uniaxial tensile tests were carried out using a 100 kN tensile frame with a
constant crosshead speed of 1 mm/min. Electric discharge machining (EDM) was used to cut
ASTM E8M [18] subsize tensile samples from the heat-treated steel panels in order to avoid any
retained austenite transformation prior to tensile testing. Furthermore, interrupted tensile tests were
conducted to determine retained austenite transformation kinetics during deformation. The retained
austenite volume fraction in all samples was measured using X-ray diffraction (XRD) analysis via
ASTM E975-13 [19], as detailed in Bhadhon et al. [11]. Mechanical properties were assessed with
respect to target 3G AHSS mechanical properties in order to select the IA parameters for the
selective oxidation and galvanizing trials.
The morphology and distribution of the external oxides formed during each stage of the thermal
processing cycle were analyzed with a JEOL 7000F field emission scanning electron microscope
(FE-SEM). Secondary electron images (SEI) of the surface were taken using an acceleration voltage
of 10 keV and a working distance of 10 mm. The depth of the solute-depleted layer formed during
the austenitization was determined from M sample cross-sections through energy dispersive X-ray
spectroscopy (EDS) elemental line scans using the JEOL 6610 SEM. An acceleration voltage of
15 keV and a working distance of 10 mm were used in this case. The external oxide thickness was
also determined during each stage of the thermal processing cycle using a Helios 5 focused ion
beam scanning electron microscope (FIB-SEM), per the method of Bhadhon and McDermid [15].
Following galvanizing trials, the coating/substrate interface of the galvanized samples was
analyzed using the JEOL 7000F FE-SEM to determine the intermetallics formed. In this regard, the
coating/substrate interface was exposed by chemically stripping both the Zn overlay and Fe-Zn
intermetallics with fuming HNO3 – leaving any interfacial Fe-Al intermetallics intact.
The coating/substrate interface cross-section of the galvanized samples was also characterized
with high-resolution scanning transmission electron microscopy (HR-STEM) using a Talos 200X in
order to identify the operative reactive wetting mechanisms. An acceleration voltage of 200 keV
was used for all samples. Site-specific TEM samples were prepared using a Helios 5 FIB-SEM. A
tungsten (W) coating was applied to the sample surface to protect the region of interest (ROI)
during FIB milling. Electron energy loss spectroscopy (EELS) was used to obtain elemental maps
of the HR-STEM sample cross-section. EELS spectra were acquired using a CMOS detector and
analyzed with Gatan Digital Micrograph software (v.3.43) to extract the elemental maps.
3. RESULTS AND DISCUSSION
3.1. Mechanical properties
Engineering stress – strain and corresponding WHR vs true strain curves of the heat-treated
med-Mn steels are shown in Fig. 2, where the global mechanical properties are summarized in
Table 3. From Fig. 2a), it can be seen that the yield strength (YS), ultimate tensile strength (UTS),
125
and total elongation (TE) were significantly affected by the IAT, where an increasing trend in UTS
and a decreasing trend in YS and TE were observed with increasing IAT. These trends are
consistent with the literature on various medium-Mn steels [2,11–13], where the authors stated that
the retained austenite chemical and mechanical stability and its resultant transformation kinetics
during deformation controlled the WHR trend and, thereby, controlled the mechanical properties of
the med-Mn steels. Fig. 2b) shows that the WHR curves in this study followed two trends,
depending on the IAT. In all cases, the WHR initially decreased and was followed by the activation
of the TRIP and/or TWIP mechanisms (marked by the green arrow in Fig. 2b)). For the M-IA
720 °C sample, the WHR decreased continuously during deformation. This WHR trend is observed
when the retained austenite is chemically and mechanically unstable and transforms rapidly during
deformation [11–13]. This was confirmed by the retained austenite transformation kinetics results
documented in Fig. 3, which lead to the conclusion that the TRIP effect was exhausted at relatively
low true strains in the case of the M-IA 720 °C sample. Contrastingly, for the M-IA 680 °C and M-
IA 700 °C samples, the WHR was sustained after the activation of the TRIP mechanism (Fig. 2b)).
This implies a relatively gradual transformation of the retained austenite owing to comparatively
higher chemical and mechanical stability, consistent with the slower retained austenite
transformation kinetics documented in Fig. 3 for these treatments. This resulted in a high UTS and
TE for both heat treatments, allowing them to achieve 3G AHSS target mechanical properties
(Table 3). Based on the results listed in Table 3, the M-IA 680 °C condition was chosen for the
selective oxidation and reactive wetting trials as it produced the highest UTS × TE product.
a)
b)
Fig. 2. a) Engineering stress vs strain and b) corresponding work hardening rate (WHR) vs true strain curves
for the heat-treated prototype medium-Mn steels.
Table 3. Mechanical properties of the as heat-treated med-Mn steels.
Sample
ID.
YS
(MPa)
UTS
(MPa)
TE
(%)
UTS × TE
(MPa%)
M-IA 680 °C
844
1130
26.4
29,800
M-IA 700 °C
740
1290
20.0
25,800
M-IA 720 °C
266
1450
13.3
19,300
126
Fig. 3. Retained austenite transformation kinetics as a function of true strain.
3.2. Oxide characterization
The 600 s austenitizing heat treatment resulted in complete coverage of the surface with nodular
oxides, as shown by the SEM micrograph in Fig. 4a). To alleviate the effects of these external
oxides on subsequent galvanizing treatments, an intermediate flash pickle was carried out prior to
the 680 °C IA. Fig. 4b) shows that most of the external oxides were dissolved by flash pickling,
resulting in exposed metallic substrate areas and extruded Fe nodules. It should be noted that the
extruded Fe nodules were the result of volume expansion associated with internal oxidation during
the austenitizing heat treatment under the –10 °C Tdp process atmosphere.
a)
b)
Fig. 4. SEM micrographs showing the surfaces of the M (–10 °C) samples a) before and b) after flash
pickling.
The austenitizing heat treatment also produced a solute-depleted layer in the near-surface region
of the steel. The elemental depth profiles obtained from the SEM-EDS line scans indicated that the
solute-depleted layer was approximately 10 μm deep (Fig. 5). This is important as Bhadhon and
McDermid [15] have shown that the solute diffusion distance during a 120 s IA treatment at 675 °C
was significantly lower than 1 µm. As a result, no significant surface change was expected after the
680 °C IA heat treatment, regardless of the IA process atmosphere Tdp, consistent with the post-IA
surface microstructures documented in Fig. 6.
These observations were further supported by the FIB trench cut analysis results shown in Fig. 7.
It can be seen that the average thickness of the external oxides was decreased significantly by the
flash pickling treatment. However, the average external oxide thickness did not change after
subsequent IA. As a result, similar to the flash pickled M(–10 °C) sample surface (Fig. 4b)), the
pre-immersion surface (i.e., after the IA treatment) comprised large area fractions of boldly exposed
metallic substrate extruded Fe nodules, and dispersed nano-scaled external oxide nodules. These
127
observations are consistent with those reported by Bhadhon and McDermid [15] and Pallisco and
McDermid [16] for their prototype med-Mn steels.
Fig. 5. SEM-EDS line scans showing elemental depth profiles for the M (–10 °C) sample.
a)
b)
Fig. 6. SEM micrographs showing the surfaces of the a) M(–10 °C)-IA 680(–30 °C) and b) M(–10 °C)-IA
680(–10 °C) samples.
Fig. 7. External oxide thickness evolution during each stage of the thermal processing cycle.
128
3.3. Galvanizing trials
Continuous hot-dip galvanizing trials following the heat treatments documented in Fig. 1 were
successful in producing high quality galvanized coatings (Fig. 8) on the prototype steel. A robust
coating with minimal bare spots was achieved regardless of the IA process atmosphere Tdp
employed. This is consistent with the results reported by Bhadhon et al. [17], where high-quality
galvanized coatings were achieved for two med-Mn steel compositions using a similar thermal
processing route.
The coating/substrate interfaces of the galvanized coatings were analyzed to evaluate successful
reactive wetting during galvanizing. SEM micrographs of the coating/substrate interfaces exposed
by fuming HNO3 are shown in Fig. 9. Significant populations of Fe-Al intermetallics (i.e.,
Fe2Al5Znx) were observed at the coating/substrate interfaces of both galvanized steels. This suggests
successful reactive wetting occurred in the 0.2 wt.% Al (dissolved) galvanizing bath.
a)
b)
Fig. 8. Uniform coating area of galvanized a) M(–10 °C)-IA 680(–30 °C) and b) M(–10 °C)-IA 680(–10 °C)
samples.
a)
b)
Fig. 9. SEM micrographs of the coating/substrate interfaces of galvanized a) M(–10 °C)-IA 680(–30 °C) and
b) M(–10 °C)-IA 680(–10 °C) samples exposed by fuming HNO3.
Reactive wetting mechanisms for the prototype steel were identified by analyzing
coating/substrate interface cross-sections via TEM-EELS. Fig. 10 shows the high angle annular
dark field (HAADF) image and TEM-EELS composite elemental map for the galvanized M(–10
°C)-IA 680(–10 °C) sample, representative of both galvanizing process atmospheres. It should be
noted that the hole in the HAADF image (Fig. 10a)) is an artifact of the FIB sample preparation
technique. Additional proof of continuous inhibition layer formation (i.e., Fe2Al5Znx intermetallics)
at the coating/substrate interface was obtained from the TEM-EELS elemental composite map (Fig.
129
10b)). Therer is evidence of “direct wetting” of the steel substrate, facilitated by the dissolution of
the exposed metallic substrate and extruded Fe nodules present on the pre-immersion surface during
galvanizing. Significant Zn ingress into the steel microstructure was also identified, as shown in
Fig. 10b). Based on Bhadhon et al. [17], it can be hypothesised that the Zn ingress into the substrate
was facilitated through two mechanisms. One was oxide cracking and spalling arising from the
differences in the coefficients of thermal expansion of the substrate and oxide species [7]. Another
was the route created by the selective dissolution of internal oxides (i.e., MnO, MnSiO3, and
Mn2SiO4) during the flash pickling step [17], leaving channels into the steel substructure.
Furthermore, Fig. 10b) shows some Fe2Al5Znx crystals were in direct contact with the oxide
nodules, suggesting “oxide wetting” was an additional reactive wetting mechanism in these
samples. However, this latter is expected to be a secondary reactive wetting mechanism owing to
the lower population of nano-scale nodular oxides on the pre-immersion substrate surface (Fig. 6).
Overall, these findings confirm that the pre-immersion surface comprising large area fractions of
boldly exposed metallic substrate, extruded Fe nodules, and dispersed nano-scale nodular oxides
was successfully galvanized owing to multiple reactive wetting mechanisms operating
simultaneously during continuous hot-dip galvanizing.
a)
b)
Fig. 10. a) HAADF image and b) corresponding TEM-EELS composite map of the galvanized M(–10 °C)-IA
680(–10 °C) sample cross-section.
4. CONCLUSIONS
The mechanical properties, selective oxidation, and reactive wetting of a prototype medium-Mn
steel (0.2C-6.25Mn-1Al-1Si-0.5Cr wt.%) with a martensitic starting microstructure were
determined for a CGL-compatible thermal processing route. Based on the results, it can be
concluded that:
1) The retained austenite formed in the 680 °C × 120 s IA sample was chemically and
mechanically stable, promoting a gradual retained austenite to martensite transformation
during deformation. This was critical in sustaining a high work hardening rate and enabled
the achievement of 3G AHSS target mechanical properties.
2) The flash pickling treatment was successful in significantly reducing the external oxide
thickness following the austenitizing heat treatment.
130
3) No significant external oxidation occurred during the IA treatment owing to the multi-
micron deep (~ 10 µm) solute-depleted layer that formed during the austenitizing heat
treatment.
4) A high-quality galvanized coating was applied to the prototype steel, which showed a well-
developed Fe2Al5ZnX inhibition layer at the coating/substrate interface, suggesting
successful reactive wetting occurred during the continuous hot-dip galvanizing. Several
reactive wetting mechanisms were identified for the prototype medium-Mn steel.
ACKNOWLEDGEMENTS
This research was financially supported by the International Zinc Association Galvanized Autobody
Partnership (IZA-GAP) program. The authors greatly acknowledge U.S. Steel Research for the
provision of the sheet steels used in this study. The authors also thank the staff of the Canadian
Centre for Electron Microscopy (CCEM), the McMaster Steel Research Centre (SRC), and the
Centre for Automotive Materials and Corrosion (CAMC) for their technical support.
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