the development of a processing window for the continuous galvanizing of a mn–cr–si martensitic...

The Development of a Processing Windowfor the Continuous Galvanizing of a Mn–Cr–SiMartensitic Steel

Anirban Chakraborty, Joseph R. McDermid,� Ranganathan Kavitha, Benoit Voyzelle,Elhachmi Essadiqi, and Frank E. Goodwin

Martensitic or complex phase steels are leading candidates for automotive impact

management applications. However, achieving high strengths while obtaining high quality

coatings via continuous galvanizing is a challenge due to cooling rate limitations of the

processing equipment and selective oxidation of alloying elements such as Cr, Mn, and Si

adversely affecting reactive wetting. The galvanizability of a Cr—Mn—Si steel with a target

tensile strength above 1250MPa was investigated within the context of the continuous

galvanizing line. The continuous cooling transformation behavior of the candidate alloy

was determined, from which intercritical and austenitic annealing thermal cycles were

developed. The evolution of substrate surface chemistry and oxide morphology during

these treatments and their subsequent effect on reactive wetting during galvanizing were

characterized. The target strength of 1250MPa was achieved and high quality coatings

produced using both intercritical (75% g) and austenitic (100% g) annealing using a

conventional 95%N2–5%H2, �308C dew point process atmosphere and 0.20wt% dissolved

(effective) Al bath, despite the presence of significant Mn and Cr oxides on the substrate

surfaces. It is proposed that complete reactive wetting by the Zn(Al, Fe) bath was

promoted by in situ aluminothermic reduction of the Mn and Cr-oxides by the dissolved

bath Al.

1. Introduction

In order to improve automotive fuel efficiency while main-

taining or improving vehicle safety, the use of advanced high

strength steels (AHSS) to reduce vehicle weight has increased

significantly in recent years and will continue to increase.[1]

Among the AHSS grades, the so-called ultra high strength

steels (UHSS) are strong candidate materials for anti-intrusion

and anti-rollover applications. The UHSS usually have a

microstructure consisting of martensite, bainitic ferrite, or

mixtures of these constituents with small amounts of ferrite.[1]

However, as in the case with all of the AHSS grades, corrosion

protection is a necessity in order to meet consumer longevity

expectations and maintain structural integrity, for which the

application of sacrificial Zn coatings by continuous galvaniz-

ing is amongst the most cost-effective solutions.

Several challenges have to be overcome in order to

process the above steels in the continuous galvanizing line

(CGL) without compromising the target tensile strength

of 1250 MPa (minimum). The challenges are as follows:

(i) obtaining the desired martensitic or complex phase

(martensiteþbainite) microstructures and (ii) providing

a substrate surface which can be reactively wet by the

Zn(Al, Fe) CGL bath to form the required h-Fe2Al5Znx

interfacial or ‘‘inhibition’’ layer and the desired metallic

Zn coating. Challenge (i) can be addressed by choosing

compositions with sufficient hardenability such that either

martensitic or martensite/bainite complex phase micro-

structures are produced using cooling rates compatible

with normal CGL operation (�308C s�1 or less). In addition,

in order to prevent martensite tempering during the 4608Chold immediately prior to and during galvanizing, the alloy

martensite start (Ms) temperature should chosen such

that it is lower than 4608C, thereby preventing martensite

formation prior to dipping into the galvanizing bath.

It is well known that in order to avoid the formation of

pearlite and promote the formation of martensite, alloying

[�] A. Chakraborty, Prof. J. R. McDermid, R. KavithaMcMaster Steel Research Centre, McMaster University, Hamilton,Ontario, CanadaEmail: [email protected]. Voyzelle, E. EssadiqiMaterials Technology Laboratory, CANMET, Ottawa, Ontario, CanadaF. E. GoodwinInternational Zinc Association, Durham, NC 27713-3210, USA

DOI: 10.1002/srin.201200124

www.steel-research.de

� 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim steel research int. 83 (2012) No. 9999 1

FULLPAPER

elements such as Mn, Al, Si, Cr, Mo, and B can be used in

steel compositions.[2–4] However, many of these alloying

elements selectively segregate to the surface during

annealing prior to galvanizing and form oxides which

can prevent reactive wetting by the CGL bath.[5–12] This

issue (challenge (ii)) can be addressed through proper

alloying element and process atmosphere oxygen partial

pressure (pO2) selection such that surfaces with sufficient

reactive wetting are formed and which are compatible with

typical CGL operation.

Some open literature is available on the galvanizing of

martensitic and complex phase UHSS, most of which are

related to electro-galvanizing or the Quench and Partition

process.[13–16] However, Song et al.[3,17] recently advocated

a C–Mn–Cr–Mo–B alloy composition which produced alloy

strengths of >1300 MPa using cooling rates of 308C s�1

during simulated galvanizing and galvannealing thermal

cycles, although it should be noted that the steels were not

actually coated and no information was provided concern-

ing the surface chemistry or microstructure following

annealing and prior to galvanizing.

The present study will examine the development of a

process window for the continuous galvanizing of a C–Mn–

Si–Cr UHSS having a target tensile strength (UTS) in excess

of 1250 MPa. In particular, the present investigation will

examine the effect of CGL-compatible thermal treatments

and process atmospheres on the microstructural develop-

ment and mechanical properties of the experimental steel

as well as the reactive wetting of the substrate surface by a

simulated CGL Zn(Al, Fe) bath. The reactive wetting of the

steel will be related to the development of the surface

chemistry and morphology as a function of the annealing

parameters and the process atmosphere composition, as

characterized by its oxygen partial pressure (pO2).

2. Experimental Methods

The chemical composition of the experimental steel is

shown in Table 1. The steel substrate was industrially

fabricated via conventional hot and cold rolling and was

supplied in the form of 1.7 mm thick cold rolled sheet with

a ferrite/pearlite microstructure. A schematic of the exper-

imental heat treatments employed throughout this study is

shown in Figure 1. Galvanizing was carried out in an iron

saturated Zn-0.2 wt% dissolved or effective Al bath main-

tained at 4608C[18] for all dipping experiments. The sample

immersion time was 4 s for all coating experiments. The

peak annealing temperature (PAT) was selected such that

75 and 100% austenite (g) were formed. The 75% g

and 100% g formation temperatures (827 and 8788C,

respectively) along with a complete continuous cooling

transformation (CCT) diagram were determined using a

BAHR-Thermoanalyse quench dilatometer. In this CCT

determination study, cooling rates for the fully austenitized

samples were varied from 0.3 to 1108C s�1. A preliminary

study on the mechanical properties was performed

on the CCT samples via microhardness assessment.

Microhardness was measured using a Clemex microhard-

ness tester using a 200 gf load coupled with a 10 s dwell

time and CMT 5.0 software. Using the results of the

dilatometry studies (discussed in detail below), a heat

treatment schedule compatible with typical CGL capabili-

ties was constructed, complete details of which are

summarized in Table 2.

All heat treatment and galvanizing experiments were

carried out using the McMaster Galvanizing Simulator

(MGS) (Iwatani Surtec). The MGS comprises an atmos-

phere controlled column equipped with a gas mixing panel

(including a resident humidification system) for atmos-

phere compositional control, a quartz lamp infrared fur-

nace used for all heat treatments and parallel platens for

controlled cooling using high flows of N2 gas. Prior to heat

treatment, samples were degreased in an 808C aqueous 2%

NaOH solution, rinsed with deionized water, cleaned ultra-

sonically in isopropanol and dried with a warm air stream.

A final cleaning with acetone was performed immediately

prior to the sample entering the galvanizing simulator. All

heat treatments were carried out under a 95%N2–5%H2

atmosphere with a controlled dew point of �308C. The

process atmosphere oxygen partial pressure as a function

of PAT is shown in Table 3, from which it can be seen that

the pO2 for the 8788C process atmosphere was approxi-

mately two orders of magnitude greater than that of the

8278C process atmosphere. The sample thermal cycle was

C Mn Si Al S P Cr

0.26 1.15 0.28 0.04 0.002 0.016 0.61

Table 1. Experimental steel chemical composition [wt%].

Figure 1. Schematic heat treatment schedule.


2 steel research int. 83 (2012) No. 9999 � 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

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controlled using a type K (0.5 mm) thermocouple welded

directly to the sample before the start of the experiment.

Experimental samples comprised 120 mm� 200 mm pan-

els with the longitudinal axis of the sample parallel to the

rolling and dipping directions. A uniform temperature and

coating area of 90 mm� 90 mm was centered in the lower

portion of the panel and only material from this area was

analyzed. A schematic of the experimental samples show-

ing the uniform coating area and control thermocouple

placement is shown in Figure 2.[19].

After heat treatment a cross-section of the substrate

was polished using conventional metallographic tech-

niques and etched using 2% nitric acid in ethanol.

Microstructural analysis of the substrate was performed

using a JEOL 7000F field emission scanning electron

microscope (FEG SEM). An acceleration voltage of

10 keV was used for all samples. The microhardness of

the as heat-treated samples were measured using the pre-

viously mentioned 200 gf load and 10 s dwell time.

Hardness measurements were taken at five different

locations for each sample.

The mechanical properties of the heat treated substrates

were assessed by tensile testing using an Instron 100 kN

electromechanical load frame with Merlin software.

Tensile samples were cut using a water jet cutter and finish

machined using carbide tools to the dimensions outlined

in ASTM E 8/E 8M-08[20] for sub-sized specimens with a

gage length of 25 mm and width of 6 mm. All tensile

samples were oriented such that the tensile axis was paral-

lel to the rolling direction. A constant crosshead speed of

1 mm min�1 was used for all tensile tests. An extensometer

was used to measure sample extension for all samples. In

all cases, five samples were tested for each annealing con-

dition and the results averaged.

In order to analyze the selective oxidation of the sub-

strate surface during annealing, small (50 mm� 10 mm)

samples were metallographically polished to remove the

effect of surface roughness. Samples were then heat

treated without galvanizing and stored in anhydrous iso-

propanol to minimize further oxidation before SEM, X-ray

photoelectron spectroscopy (XPS) and scanning Auger

microscopy (SAM) examination. Oxide morphology was

determined using SEM. SEM analysis was performed using

a JEOL 7000F field emission scanning electron microscopeFigure 2. Uniform temperature and coating area for galvanizingsimulator samples.

Sample ID. Peak annealing parameters Galvanizing parameter Cooling

Rate

[8C s�1]Temperature

(T) [8C]

Time

(t) [s]

Temperature

[8C]

Time

[s]

8278C – 60 s 827 (75% g) 60 460 4 20

8278C – 120 s 827 (75% g) 120

8788C – 60 s 878 (100% g) 60

8788C – 120 s 878 (100% g) 120

Table 2. Summary of experimental heat treatment parameters.

Peak annealing

temperature [8C]

N2 content

[vol%]

H2 content

[vol%]

Dew point

[8C]

pO2

[atm]

827 95 5 �30 8.86� 10�23

878 1.00� 10�21

Table 3. Annealing atmosphere compositions.



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and an acceleration voltage of 10 keV. Samples were sput-

ter coated with gold to avoid sample charging.

Chemical analysis of the substrate surface was carried

out using XPS. XPS analysis was performed using a PHI

Quantera X-ray photoelectron spectroscope with an Al Ka

X-ray source (Physical Electronics, Chanhassen, MN). The

spot size used was 100 mm and the take off angle was 458.All spectra were calibrated using the metallic iron peak

position at 706.62 eV. The binding energy values obtained

were accurate to within �0.1 eV and the accuracy of

the chemical compositions measurements were �5% of

the measured value in atomic percent for each element.

The binding energy results shown were collected after

sputtering with Ar to a depth of �5 nm. Elemental depth

profiles were obtained by Ar sputtering followed by XPS

analysis of the new surface. Sputtering depth measure-

ments are considered to be accurate to within �10%.

The correlation between oxide morphology and

elemental distribution was studied with elemental

mapping using SAM. SAM data were collected using a

JEOL JAMP-9500F field emission Auger microprobe. The

energy of the primary electron beam was 15 keV and

samples were tilted 308 toward the electron analyzer in

all cases. Auto probe tracking was in effect during the

collection of elemental maps to eliminate the effects of

sample drift. Auger maps were collected from the substrate

surface after Ar sputtering to a depth of 20 nm to remove

oxidation artifacts from sample exposure to ambient

atmosphere. The accuracy of the sputtering depth was

�10%.

Reactive wetting of the steels by the liquid Zn alloy bath

was studied by measuring the area of bare spots in the

galvanized coating. Measurements of any bare spots

present on the galvanized panels were performed using

a Nikon AZ100M Stereoscope on the 90 mm� 90 mm

uniform coating area. Images were captured using 10�magnification. The bare spot area was measured by man-

ually selecting the bare spot perimeter.

Reactive wetting behavior was also investigated by ana-

lyzing the steel/coating interface via SEM. An acceleration

voltage of 10 keV was used for all samples. Any reaction

products at the substrate/coating interface were exposed

for SEM analysis by stripping the zinc overlay using: (i) an

inhibited 10% H2SO4 solution in water, which leaves any

Fe—Zn intermetallics as well as any Fe–Al (Fe2Al5�xZnx

0< x< 1, i.e., h-Fe2Al5Znx) intermetallics intact and

(ii) fuming nitric acid (HNO3), which removes any Fe–Zn

intermetallics, leaving the h-Fe2Al5Znx interfacial layer

intact.

The reactions occurring at the substrate/coating inter-

face were further investigated using a JEOL 2010F trans-

mission electron microscope (TEM) equipped with an

Oxford INCA Pentafet energy dispersive X-ray (EDX) spec-

trometer (Oxford Instruments, Abingdon, Oxfordshire,

UK). The acceleration voltage was 200 keV for all samples.

EDX was performed in scanning transmission electron

microscopy (STEM) mode with a probe size of 1 nm.

TEM sample preparation was performed using focused

ion beam (FIB) milling with an NVision 40 by Zeiss.

Adherence of the coating to the substrate steel was

determined using a 1808 ‘‘U’’ bend test per ASTM A

653/A 653M-09.[21] The occurrence of any flaking of the

coating was analyzed using a Nikon AZ100M Stereoscope

at the top of the ‘‘U’’ bend sample.

3. Results

Figure 3 shows the experimentally determined CCT dia-

gram of the steel. Microhardness values for the steel as a

function of cooling rate are superimposed on the diagram.

It should be noted that the martensite start (Ms) and finish

(Mf) temperatures were determined to be �400 and 2008C,

respectively. Since the martensitic transformation takes

place after galvanizing at 4608C, it was determined there

was little possibility of tempering the martensite during

galvanizing. It was further determined that the critical

cooling rate to obtain 100% martensite was �208C s�1,

which is well suited for thermal processing in a conven-

tional CGL. It should also be pointed out that use of a

cooling rate of 108C s�1 would result in a complex phase

bainite–martensite microstructure of approximately the

same hardness (HV¼ 530) as the samples cooled at

208C s�1 (HV¼ 544). This latter finding implies that the

strength of the steel should be relatively insensitive to

cooling rates in excess of 108C s�1.

The use of a minimum annealing temperature which

will allow the target UTS of 1250 MPa (minimum) to be met

is always beneficial from the economic point of view. For

this reason, two PATs were selected based on the results of

the dilatometry studies conducted to determine the Ac1

and Ac3 temperatures (Figure 3). Based on these results,

intercritical annealing (IA) was carried out at 8278C such

Figure 3. Continuous cooling transformation diagram ofexperimental steel (F! ferrite, P! pearlite, B! bainite, andM! martensite).



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that 75% g formed and austenitic (100% g) annealing was

carried out at 8788C. As outlined in Table 2 and Figure 1,

the samples were isothermally held at their PAT for 60 and

120 s for the heat treatment trials.

Figure 4a,b show the substrate microstructures and

microhardnesses for the samples intercritically annealed

at 8278C for 60 and 120 s, respectively. Similarly, the micro-

structures and microhardnesses for the samples annealed

at 8788C for 60 and 120 s are shown in Figure 4c,d, respect-

ively. The microstructures of the samples intercritically

(75% g) annealed at 8278C contained mostly lath marten-

site along with minor amounts of ferrite (Figure 4a,b)

whereas the microstructures of the samples annealed at

8788C (100% g) comprised largely lath martensite, as

shown in Figure 4c,d. As expected, the samples annealed

at 8788C showed a higher hardness due to the small frac-

tion of ferrite in the microstructures. It should also be

observed that the 8278C – 60 s sample showed a slightly

lower microhardness (HV¼ 458) and a slightly higher vol-

ume fraction of ferrite versus the 8278C – 120 s sample

(HV¼ 484), likely a consequence of incomplete ferrite to

austenite transformation during IA in the case of the

former sample. No significant differences were found in

the microstructures and microhardnesses for the 100% g

samples.

Representative engineering stress–strain curves for the

experimental steels as a function of heat treatment sched-

ule are shown in Figure 5. A summary of the tensile results

(i.e., yield strength (YS), ultimate tensile strength (UTS),

uniform and total elongation) for the average of five trials

are provided in Table 4. From these results, it is obvious

that the minimum tensile strength of 1250 MPa was met for

all heat treatments, with the tensile strengths for a given

annealing temperature being relatively constant and

the tensile strength of the 100% g (8788C) samples being

significantly higher than those of the 75% g (8278C) inter-

critically annealed samples. This would be expected con-

sidering the increased fraction of lath martensite in the

Figure 4. SEM micrographs and steel microhardness as a function of heat treatment schedule a) 8278C – 60 s, b) 8278C – 120 s,c) 8788C – 60 s, and d) 8788C – 120 s.

0 1 2 3 4 5 6 7 8 90

200

400

600

800

1000

1200

1400

1600

1800

engi

neer

ing

stre

ss (M

Pa)

engineering strain (%)

827°C-60s 827°C-120s 878°C-60s 878°C-120s

Figure 5. Engineering stress–strain plots for the steel as a func-tion of annealing cycle.



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100% g samples versus the 75% g samples (Figure 4). An

examination of Table 4 will further show that the uniform

elongation for the samples was relatively insensitive to the

heat treatment parameters. It should be noted that a slight

increase in total elongation was observed for the 8278C –

60 s sample, likely due to the slightly higher amount of

ferrite present in this microstructure (Figure 4a).

Figure 6 shows the microstructures of the substrate

surfaces immediately prior to galvanizing as a function

of annealing condition, highlighting the substrate

surface oxide distribution. From this Figure, it can be seen

that the oxides were distributed mostly on the grain

boundaries when annealed at 8278C (Figure 6a,b) whereas

the oxides were distributed on the bulk grain surfaces

as well as grain boundaries when annealed at 8788C(Figure 6c,d).

Figure 7 and 8 show the XPS depth profiles of the as-

annealed steel surfaces as a function of annealing process

parameters. From these Figures, it can be seen that con-

siderable surface enrichment of manganese, chromium,

silicon, and oxygen was observed for all annealing

conditions. Table 5 lists the binding energies of the

elements of interest in the present study and the oxides

identified on the steel surface, in this case MnO, SiO2,

Mn2SiO4, and Cr2O3.

Scanning Auger microscopy (SAM) was used to deter-

mine the link between the as-annealed substrate surface

oxide morphology and chemistry, with the ultimate objec-

tive of linking this to reactive wetting. Figure 9 shows the

SEM micrographs of the surfaces with the corresponding

O, Fe, Mn, and Cr Auger maps for the steel annealed at

8278C (i.e, 75% g). Similarly, Figure 10 shows the SEM

Sample No. Strength Elongation

YS [MPa] UTS [MPa] Uniform [%] Total [%]

8278C – 60 s 1016 1376 5.31 8.03

8278C – 120 s 1037 1382 4.93 6.03

8788C – 60 s 1236 1588 5.07 7.65

8788C – 120 s 1201 1566 5.08 7.32

Table 4. Steel mechanical properties versus annealing conditions.

Figure 6. SEM micrographs of oxides formed on the substrate surface after annealing a) 8278C – 60 s, b) 8278C – 120 s, c) 8788C – 60 s,and d) 8788C – 120 s.



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micrographs and the corresponding O, Fe, Mn, and Cr

Auger maps for the samples annealed at 8788C (100% g).

It should be mentioned that no significant relationship

between the surface morphology and distribution of Si

was observed.

From Figure 9 and 10 it can be seen that during anneal-

ing and prior to galvanizing, considerable enrichment

of Mn and Cr in the form of oxides was observed on

all substrate surfaces for all annealing conditions.

Furthermore, the annealing temperature (and thereby

pO2) and time had a significant effect on the oxide

morphology. For the 8278C annealing temperature, it

can be seen that grain boundary oxides are more prom-

inent for the 120 s annealing time (Figure 9b) versus the

60 s annealing time (Figure 9a). Fine manganese and

chromium oxides were observed on the bulk grain surfaces

for both annealing times. As there was some correspond-

ence between the Mn, Cr, and O maps for both annealing

times, it is also possible that some nodules comprised

mixed (Mn, Cr)-oxides. As shown in Figure 10, when

annealing was carried out at 8788C the oxides of Mn and

Cr were distributed on the bulk grain surfaces as well as

grain boundaries for both annealing times. It should be

noted that the oxide morphology changed to a coarser

nodular shape when annealed at 8788C for 120 s

(Figure 10b) and that the oxides appear to cover a larger

portion of the surface versus the 8278C annealed sub-

strates. It can also be observed that there was a corre-

spondence between the Mn, Cr, and O maps for all

annealing times at 8788C, indicating the possible presence

of mixed oxides.

The ability of the zinc bath to reactively wet the steel

surface was determined by assessing the population and

size of bare spots on the coating and by observing the Fe–

Zn interface to determine if any reaction products were

present. The percentages of bare spots within the uniform

coating area as a function of heat treatment schedule are

summarized in Table 6, which clearly shows that only a

small portion of the galvanized panel comprised bare spots

with the highest portion of bare spots (1.15%) being

observed for the 8788C – 120 s case. It should also be

observed from Table 6 that the percentage of bare spots

increased with increasing annealing temperature and time.

Overall, these results indicate that the substrates were

successfully reactively wet for all annealing conditions

except for the 8788C – 120 s treatment.

Areas of the galvanized panels that demonstrated

good wettability were stripped of their zinc overlay with

0 100 200 300 400 5000

10

20

30

40

50C

once

ntra

tion

(at.%

)

Depth (nm)

O Cr Mn Fe Si

0 100 200 300 400 5000

10

20

30

40

50

Con

cent

ratio

n (a

t.%)

Depth (nm)

O Cr Mn Fe Si

a)

b)

Figure 7. XPS depth profiles of substrate surfaces after anneal-ing at a) 8278C – 60 s and b) 8278C – 120 s.

0 100 200 300 400 5000

10

20

30

40

50

Con

cent

ratio

n (a

t.%)

Depth (nm)

O Cr Mn Fe Si

0 100 200 300 400 5000

10

20

30

40

50

Depth (nm)

Con

cent

ratio

n (a

t.%)

O Cr Mn Fe Si

a)

b)

Figure 8. XPS depth profile of substrate surfaces after annealingat a) 8788C – 60 s and b) 8788C – 120 s.



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Measured elemental binding energies [eV] State (Compound)

Mn 2p3/2 Mn 2p1/2 Si 2p Cr 2p3/2 Fe 2p3/2 O 1s

641.6 653.4 103.4 576.5 706.6 530.4 MnO, SiO2, Mn2SiO4, Cr2O3,

and Fe (metallic)

Table 5. XPS identification of oxides on the substrate surface.[31–36]

Figure 9. Oxide morphology with elemental Auger maps after sputtering 20 nm a) 8278C – 60 s and b) 8278C – 120 s.



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aqueous 10% H2SO4, leaving any Fe—Zn and Fe—Al—Zn

intermetallic reaction products intact. SEM micrographs

of these samples are shown in Figure 11. Needle-like

z-FeZn13 crystals were observed for all samples as was

an underlying Fe—Al—Zn interfacial layer. A well

wetted portion of the galvanized coating area was

also stripped with fuming HNO3 leaving only any

Fe—Al—Zn intermetallics. Figure 12 presents SEM

micrographs of these samples. A continuous h-Fe2Al5Znx

layer was observed for all samples, indicating that

reactive wetting occurred for all substrates and annealing

conditions. The results of Figure 11 and 12 indicate that

the z-FeZn13 crystals likely precipitated on the surface

Figure 10. Oxide morphology with elemental Auger maps after sputtering 20 nm a) 8788C – 60 s and b) 8788C – 120 s.

Sample No. Average bare

spot size

[mm2]

Total bare

area

[mm2]

Bare

spot

[%]

8278C – 60 s 0.22 20.25 0.25

8278C – 120 s 0.29 38.88 0.48

8788C – 60 s 0.41 63.99 0.79

8788C – 120 s 0.62 93.15 1.15

Table 6. Bare spot analysis of galvanized panels (total areaanalyzed: 8100 mm2).



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of the h-Fe2Al5Znx layer during coating cooling and

solidification.

As the target mechanical properties were achieved most

economically for the 8278C – 60 s annealing condition, the

steel/coating interface was further analyzed using cross-

sectional TEM coupled with EDS analysis. Figure 13 shows

the TEM micrograph and X-ray maps for Al, Fe, Cr, O,

and Zn. From Figure 13, it can be seen that a continuous

Figure 11. SEM micrographs of the Fe–Al and Fe–Zn interfacial layer intermetallics after stripping the Zn overlay with 10% H2SO4 inwater a) 8278C – 60 s, b) 8278C – 120 s, c) 8788C – 60 s, and d) 8788C – 120 s.

Figure 12. SEM micrographs of Fe–Al interfacial layer intermetallics after stripping the Zn overlay with fuming HNO3 a) 8278C –60 s, b) 8278C – 120 s, c) 8788C – 60 s, and d) 8788C – 120 s.



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Fe—Al—Zn layer was present on this sample with a Fe–Zn

crystals being present on top of the Fe—Al—Zn layer. It

should also be noted that a Cr-oxide was present immedi-

ately below the surface of the substrate (Figure 13c–e). This

further supports the above finding of this substrate having

been successfully reactively wet by the galvanizing bath.

Coating quality and adherence were also determined by

means of a 1808 ‘‘U’’ bend test per ASTM A 653/A 653M-

09.[21] The top surfaces of the ‘‘U’’ bend test samples were

further analyzed using an optical stereoscope, the results of

which are shown in Figure 14. It can be observed that the

galvanized coating did not exhibit any flaking and showed

good adherence with the substrate. This is further evidence

that the substrates were successfully reactively wet for all

annealing conditions used in the present study.

4. Discussion

Steel composition plays a vital role in the production of

galvanized martensitic steels. Alloying elements such as

manganese and chromium improve the steel quench-

hardenability by concentrating the carbon in austenite,

which increases its stability and aids in the formation of

martensite. On the other hand, these elements also tend

to form stable oxides during annealing and can prevent

reactive wetting by the liquid zinc bath during galvanizing.

For these reasons, the composition of the steel must be

chosen in such a way that it produces sufficient quench-

hardenability to be compatible with the cooling capabili-

ties of the CGL while allowing the production of a surface

suitable for reactive wetting by the CGL bath.

Figure 3 through Figure 5 and Table 4 show that the

substrate composition and all experimental heat treat-

ments used (Table 2) met the requirements of producing

the desired microstructures and mechanical properties

(UTS�1250 MPa). Indeed, the mechanical property results

(Table 4 and Figure 5) show that the mechanical property

target was significantly exceeded for all experimental heat

treatments used. As would be expected, the increased

volume fraction of austenite, and consequently martensite,

in the steels annealed at 8788C (Figure 4) resulted in a

higher tensile strength being realized without a significant

deterioration in uniform elongation. This would be

expected due to the higher volume fraction of martensite

in the fully austenitized versus intercritically annealed

microstructures.[22]

Figure 6 through Figure 10 and Table 5 summarize

the experimental findings with respect to the surface struc-

ture and chemistry prior to the substrates undergoing

reactive wetting by the Zn(Al, Fe) bath. Both Figure 6

and the elemental Auger maps in Figure 9 and 10 show

that the oxide morphology comprised grain boundary films

Figure 13. TEM micrograph and X-ray maps across the substrate/coating interface for the 8278C – 60 s sample a) bright field image,b) Al map, c) Fe map, c) Cr map, e) O map, and f) Zn map.



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with an increasing area of the bulk grains being covered by

nodular oxides as the annealing time and temperature

increased. These nodules comprised mainly manganese

(Mn) and chromium (Cr) oxides with some evidence of

mixed (Mn, Cr) oxides being present on the bulk grain

surfaces and grain boundaries. This overall trend was con-

firmed by the XPS profiles (Figure 7 and 8), which showed

considerable enrichment of manganese (as MnO, Table 5)

and chromium (as Cr2O3) with minor enrichment of silicon

(as SiO2). For example, the manganese enrichment was

�20 at% when annealed at 8278C (Figure 7) versus�25 at%

when annealing at 8788C (Figure 8). Similar trends are also

seen for the chromium enrichment, where there was

�5 at% Cr at the surface when annealing at 8278C whereas

there was �10 at% Cr on the surfaces annealed at 8788C.

These observations are not surprising given that the selec-

tive oxidation process is diffusion driven and that the

atmospheric pO2 is two orders of magnitude greater for

the 8788C annealing temperature (Table 3), thereby pro-

viding a more significant oxygen flux into the substrate. It

should also be noted that internal oxidation of chromium

was clearly observed for all annealing conditions, where

the peak chromium concentration was invariably located

some 25 nm below the initial surface. It should be further

noted that internal oxidation of chromium was observed

for the cross-sectional TEM study conducted on the

galvanized 8278C – 60 s sample (Figure 13a,d,e), supporting

this finding. From these observations, it can be concluded

that the post-annealing surfaces comprised mixtures of Mn

and Cr-oxide nodules and grain boundary films whose

thickness and coverage increased with both annealing

temperature (and thereby increase atmospheric pO2)

and annealing time, with some internal oxidation of Cr

having taken place.

The interaction of the as-annealed surfaces with the

0.20 wt% dissolved Al galvanizing bath are highlighted

in Figure 11 through Figure 14 and Table 6. From the

interfacial microstructures shown in Figure 11 and 12 com-

bined with the bare spot analysis shown in Table 6, it is

obvious that a high degree of reactive wetting was obtained

for all annealing conditions except for the 8788C – 120 s

thermal cycle. In addition, Figure 14 shows that all of the

coatings produced showed high adherence, indicating that

the Fe–Zn phases shown in Figure 11 did not have a del-

eterious effect on the coating properties. From the above, it

can be concluded that, despite the presence of significant

oxides on the as-annealed surfaces, the substrate surfaces

annealed at 8278C and 8788C – 60 s were reactively wet and

formed the desired Fe–Al interfacial layer to produce high

quality, adherent metallic Zn coatings. The exception to

this was the high temperature, longer annealing time 8788C– 120 s condition. However, since there are no significant

strength benefits to be gained by using this annealing cycle

versus the 8788C – 60 s cycle, there is no practical reason

to employ this cycle in processing these substrates

industrially.

From Figure 11 through Figure 13, it can be seen that

a well developed Fe–Al interfacial layer was formed

for the majority of the annealing conditions explored

and in particular for the lower energy intensity annealing

Figure 14. Stereo micrographs from the top of the 1808 ‘‘U’’ bend galvanized steel samples a) 8278C – 60 s, b) 8278C – 120 s,c) 8788C – 60 s, and d) 8788C – 120 s.



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temperature of 8278C (Figure 12). From the XPS and Auger

data, which show the presence of significant population

and thicknesses of oxides in the form of nodules on the

substrate surfaces immediately before dipping, this would

imply that a mechanism exists by which the oxide layer was

either removed as a barrier for reactive wetting. This oxide

removal mechanism exists was confirmed by the TEM

observations of Figure 13, where no visible oxide phases

were observed at the coating/Fe–Al layer interface in the

case of the 8278C – 60 s sample.

Among the more promising candidates for this oxide

removal mechanism is the so-called in situ alumino-

thermic reduction mechanism proposed by Khondker

et al. [23] for MnO on the surface of Mn-containing dual

phase steels by the dissolved Al in continuous galvanizing

baths. The proposed reaction was recently confirmed by

Kavitha and McDermid,[24] who also showed that a 100 nm

monolithic MnO layer could be reduced by a 0.20 wt% Al

(dissolved) galvanizing bath in the 4 s immersion times

employed in the present study. The chemical reaction

occurring during this process can be written as:[23,24]:

2½Al� þ 3MnO �!460CAl2O3 þ 3½Mn� (1)

where the square brackets denote dissolved species in the

liquid Zn-alloy bath and compounds without brackets

indicate pure condensed species. For the MnO thicknesses

observed on the present substrates (�100 nm or less thick

for all but the 8788C – 120 s sample), the MnO reduction

kinetics shown by Kavitha and McDermid[24] indicate

that this layer would be reduced in the 4 s immersion

times employed in the present study. Furthermore, since

the surface oxide microstructures and Auger maps

in Figure 9 and 10 indicate that the substrate surfaces

generally comprised thin oxide layers with more localized

nodules, the aluminothermic reduction mechanism clearly

had sufficient kinetics to produce the full reactive wetting

and low bare spot counts observed for the substrates

annealed at 8278C (Figure 12a,b, Table 6) and the 8788C– 60 s sample (Figure 12c, Table 6). In the case of the 8788C– 120 s, the thicker MnO layer observed in the XPS analysis

(�200 nm) could not necessarily be completely reduced by

the resident dissolved Al,[24] resulting in the increased bare

spots (Table 6) and holes in the Fe-Al interfacial layer

(Figure 12d) observed for this sample.

Per the XPS and elemental Auger plots in Figure 7

through Figure 10, considerable Cr and/or mixed (Mn,

Cr)-oxides were also observed on the substrate surfaces.

The evidence from the Fe–Al interfacial layer formation

(Figure 12) also implies that a mechanism for removal of

these oxides in the galvanizing bath was also present. It is

thus proposed that aluminothermic reduction of Cr2O3

also occurred at the interface during substrate immersion

in the Zn—Al—Fe bath:

2½Al� þ Cr2O3 �!460C

Al2O3 þ 2½Cr� (2)

In proposing reaction (2), it should be noted that the free

energy of formation of Cr2O3 is somewhat less negative

than that of MnO,[25] further implying that the free energy

of formation of the (Mn, Cr) mixed oxides are intermediate

to those of Cr2O3 and MnO. This in turn implies that

the aluminothermic reduction of the Cr and (Cr, Mn)-

oxides is more energetically favorable than that of the

MnO reduction reaction.

Further evidence for the consumption of Al at the

substrate bath interface per reactions (1) and (2) can be

found in Figure 11, which shows the Fe–Zn and Fe—Al

intermetallics at the substrate–coating interface. Per

Figure 12 and 13, the Fe–Zn layer precipitated on the

Fe–Al layer during coating solidification. It has been estab-

lished by several authors[26–30] that the Fe—Al interfacial

layer grows outward into the liquid Zn coating by Fe dif-

fusion through the fine-grained inner h-Fe2Al5Znx layer. In

the present case, the Fe diffusing through the h-Fe2Al5Znx

layer met a mass transport layer considerably depleted in Al

and supersaturated with Fe, where the precipitation of Fe–

Zn intermetallics was preferred.[18] It should be pointed out

that the thicker the oxide film (e.g., the 8278C samples versus

the 8788C – 120 s sample) and the greater the Al consumption

required for reactions (1) and (2), the greater the population

of Fe–Zn intermetallics at the substrate–coating interface

(Figure 11), indirectly supporting this theory.

The overall conclusions of this study are that a Fe–

1.2Mn–0.3Si–0.6Cr–0.26C martensitic steel with a tensile

strength exceeding 1250 MPa can be produced using proc-

ess atmospheres, cooling rates and galvanizing bath com-

positions compatible with the capabilities of conventional

CGLs. In this case, the minimum energy route to produce

suitable coatings and mechanical properties was described

using the 75% g intercritical 8278C – 60 s annealing cycle,

although satisfactory coatings and higher tensile strengths

could be obtained using the 100% g 8788C – 60 s thermal

cycle. The thermal cycles employed resulted in the selec-

tive surface oxidation of Mn and Cr, but which did not

inhibit full reactive wetting due to the proposed alumino-

thermic reduction of the Mn and Cr-oxides by the resident

bath dissolved Al.

5. Conclusions

The hot dip galvanizability of an ultra high strength

(UTS� 1250 MPa) martensitic steel was systematically

analyzed to determine a suitable process window for pro-

duction in conventional CGLs. From this study, the follow-

ing conclusions can be drawn:

(1) The target tensile strength of 1250 MPa (minimum) was

achieved using both intercritical (75% g) and austenitic

(100% g) annealing for 60 and 120 s followed by cooling

to room temperature at 208C s�1. However, the target



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strength was achieved most economically when

employing the 8278C – 60 s annealing cycle.

(2) Surfaces which could be successfully reactively wet by

the Zn (Al, Fe) bath were produced using a conventional

95%N2–5%H2, �308C dew point process atmosphere for

the intercritically annealed (8278C – 60s, 8278C – 120s)

and austenitically annealed sample 8788C – 60s, despite

the presence of significant thicknesses (100 nm) of Mn

and Cr-oxides on the bulk grain surfaces and in the form

of grain boundary films. Reactive wetting was adversely

affected in the case of the 8788C – 120s sample due the

presence of a thicker oxide layer on this sample.

(3) The morphology of the surface oxides changed when

annealing was carried out at higher temperatures and

the oxide layer thickness increased with annealing

temperature (and thereby process atmosphere pO2)

and annealing time on the surface of the austenitically

annealed steel. It was also observed that these oxides

were distributed mostly on the grain boundaries when

annealed at 8278C whereas they were distributed on the

bulk grain surface as well as grain boundaries when

annealed at 8788C.

(4) It is proposed that both the Mn and Cr-oxides on the

surface of the steel were reduced and reactive wetting

promoted due to in situ aluminothermic reduction of

MnO and Cr2O3 by the dissolved bath Al. This mech-

anism was not able to fully reduce the oxide layers on

the 8788C – 120s due to the thicker oxides observed on

this sample.

(5) A Fe–1.2Mn–0.3Si–0.6Cr–0.26C martensitic steel with a

tensile strength significantly exceeding 1250 MPa was

produced using process atmospheres, cooling rates

and galvanizing bath compositions compatible with

the capabilities of conventional CGLs. In this case,

the minimum energy route to produce suitable coat-

ings and mechanical properties was described using

the 75% g intercritical 8278C – 60s annealing cycle.

Acknowledgments

The authors would like to thank the members of the

International Zinc Association, Galvanized Autobody

Partnership, ArcelorMittal (Dofasco), and the Natural

Sciences and Engineering Research Council of Canada

(NSERC) for their financial support and provision of exper-

imental materials. The author would like to thank the

RIEM program of CANMET-MTL for access to the quench

dilatometer for determination of the CCT diagram.

The authors would also like to thank John Thomson,

Mariana Budiman, Doug Culley, Steve Koprich and Rob

Lemmon for their technical support.

Received: May 18, 2012

Keywords: aluminothermic reduction; continuous

cooling transformation; continuous galvanizing;

martensitic steel; reactive wetting

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the development of a processing window for the continuous galvanizing of a mn–cr–si martensitic...

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