1

Hydrogen permeation in dual phase (DP) and quenched and partitioned

(Q&P) advanced high strength steels (AHSS) under cathodic charging and

simulated service conditions

Qinglong Liu

1,*, Jeffrey Venezuela

1, Mingxing Zhang

1, Qingjun Zhou

2, Andrej Atrens

1

1 The University of Queensland, Division of Materials, School of Mining and Mechanical

Engineering, St. Lucia, 4072 Australia * Corresponding author, andrejs.atrens@uq.edu.au

2 Baoshan Iron & Steel Co., Ltd, Research Institute, Shanghai, 201900, China

Abstract: Hydrogen permeation through DP and Q&P advanced high strength steels (AHSS) was

investigated (i) for cathodic charging at different potentials in 0.1 M NaOH, (ii) for immersion in 3wt%

NaCl solution at the free potential simulating the corrosion of steel in service and at zinc potential for

corrosion of the galvanized steel during service, and (iii) for immersion in 0.1 M HCl solution at the

free potential. For cathodic charging, hydrogen permeation and hydrogen concentration decreased

with less negative charging potentials. Hydrogen permeation and hydrogen concentration under

simulated service conditions in 3wt% NaCl were lower than those under the least negative cathodic

charging potential of -1.100 VHg/HgO in 0.1 M NaOH. For galvanized steels, once the Zn coating was

corroded, more hydrogen would be introduced into and permeate through the steel. The hydrogen

concentration was also measured by a hot extraction analyser, providing consistent results with those

from the permeability experiments. Key words: steel; hydrogen permeation; advanced high strength steel

1. Introduction

In recent decades, advanced high strength steels (AHSS) were developed, and have

been adopted, for auto components to meet the demands for light weight and increased

vehicle safety [1, 2]. Among AHSS, DP and TRIP steels exhibit a good combination of

strength and ductility, and good energy absorption ability. Quenching and partitioning (Q&P)

is a new heat treatment to produce TRIP steel. This process was designed with the desire for

producing steels with a combination of better strength and ductility, and high strain hardening

capacity, achieved by a microstructure containing more retained austenite and a greater TRIP

effect. Q&P steels are thus third generation AHSS with high strength and high ductility [3].

Hydrogen can be introduced into the steels, during steel making, or during auto

construction processes such as painting, or during corrosion practical service. The hydrogen

evolution reactions are as follows [4]:

H3O+ + M + e MHads + H2O (acid) (1)

H2O + M + e MHads+ OH- (neutral, alkaline) (2)

where M represents the metal surface, and MHads represents the adsorbed hydrogen on the

metal surface. The atomic hydrogen may be desorbed by the desorption Eq. (3), through

which two adsorbed hydrogen atoms combine to form a molecule of hydrogen that leaves the

metal surface, or by the electrochemical desorption reaction given by Eq. (4) or (5).

2MHads →H2 + 2M (3)

MHads + H3O+ + e H2 + H2O + M (acid) (4)

MHads + H2O + e H2 + OH- + M (neutral, alkaline) (5)

Some of the adsorbed hydrogen enters the metal, MHabs, by the following equilibrium

reaction:

2

MHads ↔ MHabs (6)

This hydrogen, in combination with an applied stress, interacts with the steel and can

cause the degradation of the mechanical properties of the steel and even catastrophic failure.

This phenomenon of hydrogen embrittlement (HE) and is a possible concern for AHSS [5, 6].

Thus it is important to understand how much hydrogen can enter into and permeate through

the AHSS in various service environments. This research is part of a wider research program

that seeks to understand the influence of hydrogen on the mechanical behaviour of AHSS,

and to provide information for their applications in auto industry.

The permeability experiments based on the method of Devanathan and Stachurski [7]

have been widely used for studying hydrogen transportation in metals.

This paper builds on existing research that studied hydrogen embrittlement (HE) of

AHSS under increasingly severe hydrogen charging conditions [8-10], from inert, moderate,

to severe, and summarises our prior studies on hydrogen permeation in DP and Q&P steels

under various conditions [11, 12], including (i) cathodic charging in alkaline solution, (ii)

simulated service corrosion and (iii) free corrosion in 0.1 M HCl solution.

2. Experimental

2. 1. Materials

The DP and Q&P steels were rolled sheets, and were designated as 980 DP, 1200 DP-

GI and 980 QP. “GI” referred to galvanization of the steel, which involving applying a Zn

corrosion-protection coating on the steel. The average as-received sheet thicknesses were 1.35

m for 980 DP and 1200 DP-GI, and 1.93 mm for 980 QP. The specimens were cut to 30 × 30

mm with the as-received thickness. The chemical composition, mechanical properties and

microstructural composition of the steels are presented in Table 1.

Table 1 Chemical composition (in wt %), mechanical properties and microstructural

composition of the studied steels. Steel

designation

C Si Mn Yield

stress,

MPa

Tensile

stress,

MPa

Elongation

at fracture,

ef, %

ferrite

%

martensite

or banite

%

retained

austenite

%

980 DP 0.085 0.276 2.255 592 930 8.3 40 60 0

1200 DP-GI 0.121 0.244 2.425 896 1198 4.7 26 74 0

980 QP 0.209 1.386 1.876 682 1020 11.3 39 53 8

2.2. Permeability experiments

(i) Cathodic charging

A double-cell arrangement was used for permeability experiments based on that of

Devanathan and Stachurski [7]. Each cell contained 0.1 M NaOH solution, and a three

electrode system with the steel as the working electrode, a Pt wire as the counter electrode,

and an Hg/HgO, KOH (20%) reference electrode connected to a Luggin capillary.

The 980 DP, 1200 DP-GI and 980 QP specimens were ground to: 0.53, 0.70 and 0.48

mm, respectively. The zinc coatings on the top and bottom surfaces of the galvanized 1200

DP-GI steel were ground off. The area exposed to the solution on the hydrogen-exit side was

3.394, 2.997 and 2.930 cm2, respectively. The specimen side exposed to hydrogen charging

cell was polished to 3 µm, washed with distilled water and ethanol, and dried. The hydrogen-

exit side of the specimen was plated with palladium to prevent oxidation of the steel.

Hydrogen was produced by a negative potential applied with a MP 81 potentiostat. A

PARSTAT 2273 maintained a potential of + 300 mVHg/HgO to oxidise the emerging hydrogen

and measure the amount of hydrogen that permeated through. N2 was bubbled throughout the

whole experiment to remove oxygen, which could contribute to the oxidation current density.

The exit side background current density was lower than 0.2 μA cm-2

before cathodic

3

charging on the hydrogen-entry side. This background current was subtracted from the

permeation current.

An uninterrupted pre-charging at -1.600 VHg/HgO was conducted for 60 h on the

hydrogen-entry side of the specimen in order to stabilise the surface. During pre-charging

there was typically a steady increase in permeation current density until a steady state was

reached. Thereafter, successive transients were measured, such as those from -1.600 VHg/HgO

to -1.700 VHg/HgO, as well as a transient loop from -1.700 VHg/HgO to -1.100 VHg/HgO and back

to -1.700 VHg/HgO. The potentials for the transient loop were: -1.700 VHg/HgO, -1.600 VHg/HgO,

-1.400 VHg/HgO, -1.200 VHg/HgO, -1.100 VHg/HgO, and similarly back to -1.700 VHg/HgO. The

same experimental sequence was carried out for each steel. All the experiments were carried

out at room temperature, 23 ± 2 ℃ .

The permeation transients can be expressed by:

in = 𝑖p− 𝑖p

0

𝑖p∞− 𝑖p

0 = 2𝐿

√𝜋𝐷𝑡∑ exp( −

(2𝑛+1)2𝐿2

4𝐷𝑡)∞

𝑛=0 (Rise transients) (7)

in = 𝑖p− 𝑖p

∞

𝑖p0− 𝑖p

∞ = 1−2𝐿

√𝜋𝐷𝑡∑ exp( −

(2𝑛+1)2𝐿2

4𝐷𝑡)∞

𝑛=0 (Decay transients) (8)

where in is the normalized current, ip is the measured permeation current density at time t, 𝑖p0

is the initial steady-state permeation rate at time t = 0, 𝑖p∞ is the new steady-state permeation

current density, and L is the thickness of the steel sheet. For the first charging, 𝑖p0 = 0, and for

the complete decay 𝑖p∞= 0. The experimental permeation curves were fitted to Eq. (7) or (8),

to determine the effective diffusion coefficient, Deff, and then the total hydrogen

concentration, CH, at the sub-surface on the cathodic side, can be calculated from:

CH = 𝑖𝑝

∞ 𝐿

𝐹 𝐷eff (9)

where F is the Faraday constant, L is the thickness of the membrane, 𝑖p∞ is the steady-state

permeation rate, and Deff is the measured effective diffusion coefficient of hydrogen in the

steels.

(ii) Simulated service corrosion

These permeability experiments used the same set-up and experiment conditions with

3wt% NaCl solution with pH 7.3 on the entry side either (i) with the steel at the free corrosion

potential, or (ii) with the steel surface polarised to the free corrosion potential of zinc, in order

to simulate the maximum hydrogen concentration for steel corroding in contact with a

galvanised layer.

For the experiments at the free potential in 3wt% NaCl solution, the 980 DP, 1200

DP-GI and 980 QP specimens were ground to 0.71, 0.86 and 0.47 mm, respectively. The area

exposed to the solution in the hydrogen-exit-side cell was 3.309, 3.423 and 2.588 cm2,

respectively. The open circuit potential was measured to be -0.668 VAg/AgCl for the steels

freely corroded in the 3wt% NaCl solution. The free corrosion potential was constant and did

not change significantly with time.

The experiments at the Zn potential in 3wt% NaCl solution used an applied potential

of -0.950 VAg/AgCl, which was the measured open circuit potential of the Zn coating in 3wt%

NaCl solution. The reference electrode was Ag/AgCl, KCl (saturated) and the counter

electrode was a Pt wire. The 980 DP, 1200 DP-GI and 980 QP specimens were ground to

0.56, 0.90 and 0.55 mm, respectively, and exposed an area of 2.608, 2.912 and 3.243 cm2,

respectively, to the deaerated 0.1 M NaOH solution on the hydrogen-exit side.

(iii) 0.1 M HCl

Experiments at the free potential in 0.1 M HCl solution, which was much more severe

than any environment encountered in auto service, were carried out using 980 DP steel. These

experiments were carried out because this solution has been used to study the influence of

4

hydrogen on AHSS [13, 14]. The specimen was ground to 0.91 mm and exposed area of

2.608 cm2 in the hydrogen-exit-side cell. The open circuit potential was measured to be

-0.476 VAg/AgCl.

2.3. Diffusible hydrogen concentration

The diffusible hydrogen concentration was measured using a BRUKER G4

PHOENIX DH carrier gas hot extraction analyser for 980 DP, 1200 DP-GI and 980 QP steels

which were uncharged as well as were (i) under hydrogen chathodic charging in 0.1 M NaOH

solution at over-potential values of -0.257 V, -0.557 V and -0.857 V, (ii) at free corrosion in

3wt% NaCl solution, simulating free corrosion of the steel of the car body in service, (iii)

cathodically polarised to the zinc potential in 3wt% NaCl to simulate the maximum amount

of hydrogen produced for the steels protected by a zinc coating applied by galvanisation, and

(iv) at free corrosion potential in 0.1M HCl, all for 24 h. The sample dimensions were

10 mm × 50 mm with the as-received thicknesses. All sample surfaces were ground to 1200

grit using SiC paper.

After 24-hour hydrogen charging, the sample was cleaned with distilled water and

ethanol, dried with blowing air, weighed, inserted in the quartz extraction tube in the

temperature-programmable infrared heated furnace, and the furnace was heated by the

instrument to 400 °C, which was used because it was expected that all diffusible hydrogen

would be released but that irreversibly bound hydrogen would not be released. The diffusible

hydrogen was released into the carrier gas flow of pure nitrogen, and the amount of hydrogen

was determined from the increase in thermal conductivity of the carrier gas. There was a

constant time interval of 3 min between the end of hydrogen charging and the beginning of

the hydrogen measurement.

3. Results

3.1. Cathodic charging

Fig. 1(a) presents the hydrogen permeation current density versus time for the three

steels during the 60-h pre-charging in the 0.1 M NaOH solution at -1.600 VHg/HgO, which was

started after the background current density in the right-hand cell decreased to less than

0.2 μA cm-2

. Thereafter, successive transients as well as a transient loop were measured.

The permeation current density increased significantly after the break through, and

then reached a maximum value, as shown in Fig. 1(a). The maximum current density

indicated approximately steady-state hydrogen charging conditions, and was 32 μA cm-2

,

21 μA cm-2

and 7 μA cm-2

, for 980 DP, 1200 DP-GI and 980 QP steel, respectively. The

significant increase in permeation current density with cathodic pre-charging time was

consistent with the literature [9, 15], and is attributed to a stabilisation of the steel surface due

to (i) the reduction of some of the air-formed oxide, which increased the surface coverage of

hydrogen, and (ii) the weakening of the bonding force between the adsorbed hydrogen and

the steel, which facilitated hydrogen absorption.

Fig. 1(b) shows a typical transient loop from -1.700 VHg/HgO to -1.100 VHg/HgO and

back to -1.700 VHg/HgO in 0.1 M NaOH for 980 QP steel. Each transient was a partial transient,

which meant that there was always a significant charging potential applied to the specimen,

and the specimen was not allowed to completely discharge all hydrogen. Similar transient

loops were also carried out for the other DP and Q&P steel grades. The experimental

permeation transients, such as those from the transient loop shown in Fig. 1(b), were fitted to

Eqs. (7) and (8) by Matlab to determine the hydrogen diffusion coefficient values, Deff, and

the total hydrogen concentration, CH, at the sub-surface on the cathodic side, was calculated

by Eq. (9).

5

Fig. 1 (a) Hydrogen permeation curves for the steels during cathodic pre-charging at-1.600

VHg/HgO in NaOH solution and (b) hydrogen permeation transients at different cathodic

charging potentials for the 980 QP steel after the pre-charging (P1: - 1.700 VHg/HgO, P2:

- 1.600 VHg/HgO, P3: - 1.400 VHg/HgO, p4:- 1.200 VHg/HgO and P5:- 1.100 VHg/HgO).

For the DP steels studied in our research, the values for effective diffusion coefficient,

as were presented in our prior study [11], were in the range of 0.7 × 10-6

~ 2.1 × 10-6

cm2 s

-1,

consistent with those from literatures [16-19]. For 980 QP steel, the effective diffusion

coefficient varied from 3 × 10-7

cm2 s

-1 to 7 × 10

-7 cm

2 s

-1, in agreement with the range of

0.3 × 10-7

cm2 s

-1 to 6 × 10

-7 cm

2 s

-1 from other studies [17, 20, 21], regardless that different

methods were used to determine the diffusion coefficient, which would lead to a variance of

the values. For instance, Yang et al. [20] used two methods, the time lag method and the

breakthrough method, to determine the hydrogen diffusion coefficient of their Q&P steels.

The results obtained from both methods had a difference of least a factor of 2. In addition, the

values of the diffusion coefficient were similar for the first two rise transients from -1.600

VHg/HgO and -1.700 VHg/HgO after long time pre-charging, and the fitting of the permeation

curve to the equations were good, as indicated in our prior research [11, 12]. These values of

the diffusion coefficient were identified as the lattice diffusion coefficient, DL. The diffusion

coefficients measured for the subsequent decay transients gave decreasing values of Deff.

Thereafter the partial rise transients gave values of the effective diffusion coefficient, Deff,

which increased towards the value for the lattice diffusion coefficient.

Fig. 1(b) indicates that as the charging potential was changed from -1.700 VHg/HgO to

-1.100 VHg/HgO the steady-state current density and the hydrogen concentration at each

potential decreased, indicating less hydrogen entered and permeated through the steel at a less

negative charging potential. When the charging potential was increased from -1.100 VHg/HgO

to -1.700 VHg/HgO, there was more hydrogen entering and permeating through the specimen,

providing a higher steady-state current density and hydrogen concentration. This trend was

consistent with those from Liu et al. [9] and Venezuela et al. [10], where they found that

during cathodic charging in 0.1 M NaOH solution, an increasingly negative charging

potential lead to an increased hydrogen fugacity, and thus an increased hydrogen

concentration and permeation current density in the steel.

Fig. 2 presents the experimentally determined relationship between the total hydrogen

concentration at the input side of the permeation specimen, CH, and the over-potential, η, for

all the steels under various cathodic hydrogen charging potentials in the 0.1 M NaOH solution.

For clarity, fitting lines are shown only for 1200 DP-GI and 980 QP. There was a turning

point for each steel at an over-potential of about -0.5 V. The relationship between ln CH and η

was linear both below and above the turning point, similar to that found in other research [9,

22].

6

Fig. 2 The hydrogen concentration, CH vs. over-potential, η, for the studied steels for various

conditions.

3.2. Simulated service corrosion

Fig. 3(a) presents the hydrogen permeation current density versus time for all the

steels under simulated auto service conditions in 3wt% NaCl solutions at Ecorr and EZn. These

are not permeation transients, as is apparent by a comparison with Fig. 1.

Fig.3 Hydrogen permeation current density versus time (a) for the steels for simulated service

corrosion and (b) for 980 DP steel under simulated service corrosion as well as under free

corrosion in 0.1 M HCl solution.

Fig. 3(a) shows that for the steels charged at EZn, the permeation curves experienced a

fast rising at the beginning to reach a maximum, then declined somewhat after a few hours,

attributed to the formation of an oxide film on the cathodic polarized surface, and finally

slowly increased with time. The permeation current density reached a steady state after about

60-hour charging. The steady-state permeation current density could be higher than the

maximum value during the fast rising stage. For the steels charged at Ecorr, the curves all

consisted of two stages: a rising stage and a steady state stage. The increase in the permeation

current density over this long time is attributed to changes at the ingress side of the specimen.

The hydrogen permeation current density increased with time, due to an increased amount of

hydrogen atoms that formed and permeated through the steel during the corrosion reaction,

and reached a steady state when the ingress surface conditions had reached steady state [9].

After about 60 hour at Ecorr in 3wt% NaCl solution, most of the surface area of the studied

steels was covered by a brown corrosion layer, which easily chipped off after removal of the

specimen from the cell, revealing a black layer of magnetite [23].

7

Table 2 presents the values of (i) the steady state permeation current density, i∞, (ii)

the effective hydrogen diffusivity, Deff, and (iii) the total hydrogen concentration, CH, under

simulated service conditions.

Table 2 Permeability parameter values for the studied DP and QP steels under (i) simulated

service conditions in 3wt% NaCl solution and (ii) at free corrosion in 0.1 M HCl solution. Steel

designation

Condition Over-potential

(η, V)

i∞

(µA cm-2

)

Deff

(cm2 s

-1)

CH

(µg g-1

)

Cd

(µg g-1

)

980 DP

at Ecorr in 3wt% NaCl -0.039 0.095 4.39 × 10-7

0.020 0.031

at EZn in 3wt% NaCl -0.322 1.480 8.78 × 10-7

0.133 -

at Ecorr in 0.1M HCl -0.203 0.748 7.24 × 10-7

0.129 -

1200 DP-GI at Ecorr in 3wt% NaCl -0.039 0.090 4.62 × 10

-7 0.022 0.053

at EZn in 3wt% NaCl -0.322 0.582 7.03 × 10-7

0.098 0.090

980 QP at Ecorr in 3wt% NaCl -0.039 0.053 1.71 × 10-7

0.018 0.010

at EZn in 3wt% NaCl -0.322 0.438 3.03 × 10-7

0.109 0.100

Fig. 3(a) and Table 2 show that the steady state permeation current density, i∞,

decreased in the following order for the steels at both EZn and Ecorr: 980 DP, 1200 DP-GI, and

980 QP. In addition, the i∞ values were significantly higher when hydrogen was charged at

EZn than for those at Ecorr in the 3wt% NaCl solution, as was also the case for the values of the

total hydrogen concentration, CH. For example, the values of i∞ and CH for 980 DP steel were

0.020 µg g-1

and 0.095 µA cm-2

at Ecorr, and 0.133 µg g-1

and 1.480 µA cm-2

at EZn, increased

at least by about 7 times.

Fig. 2 also presents the total hydrogen concentration at the input side of the steel

specimen, CH, for all the steels under simulated service corrosion, and provides a comparison

with the hydrogen concentrations produced by cathodic hydrogen charging.

3.3. In 0.1 M HCl solution

Fig. 3(b) presents the permeation current density versus time curves for 980 DP steel

at free corrosion in 0.1 M HCl solution and compares the data with those measured

previously under simulated auto service corrosion in 3wt% NaCl solutions at Ecorr and EZn, as

in Fig. 3(a). The permeation current density increased faster than that in 3wt% NaCl solutions

at EZn, reaching to a nearly steady-state after about 5 hours and remained steady thereafter.

Values of the steady state permeation current density, i∞, and the total hydrogen concentration,

CH, for 980 DP steel at Ecorr in 0.1 M HCl solution, were also included in Fig. 2 and Table 2,

and were found to be significantly higher than those for 980 DP at Ecorr in 3wt% NaCl

solution. The steel surface was observed to be fully covered by corrosion products in the 0.1

M HCl solution. The corroded area was measured after the experiments to be 3.002 cm2

and

2.014 cm2

for the steel at Ecorr in 0.1 M HCl and in the 3wt% NaCl solution, respectively.

3.4 Diffusible hydrogen content

Table 3 and Fig. 4 present the diffusible hydrogen concentration, Cd, for the studied

steels for (i) being uncharged, (ii) hydrogen cathodic charging in 0.1 M NaOH solution at

over-potential values of -0.257 V, -0.557 V and -0.857 V, (iii) free corrosion in 3wt% NaCl

solution, simulating free corrosion of the steel of the car body in service, (iv) cathodic

polarisation of the steel to the zinc potential in 3wt% NaCl to simulate the maximum amount

of hydrogen produced for the steels protected by a zinc coating applied by galvanisation, and

(iv) free corrosion in 0.1 M HCl solution.

The values of diffusible hydrogen concentration, Cd, for the uncharged samples were

not higher than 0.002 µg g-1

and were significantly lower than those at Ecorr in 3wt% NaCl.

The Cd values at Ecorr in 0.1 M HCl were much higher than those at Ecorr in NaCl. Comparing

the simulated auto service conditions in NaCl, the hydrogen concentrations at EZn were higher

than the corresponded values at Ecorr. For example, the diffusible hydrogen concentration for

8

980 QP at EZn in 3wt% NaCl was 0.100 µg g-1

, 10 times higher than 0.010 µg g-1

at Ecorr. For

those under cathodic charging potentials in NaOH solution, the hydrogen concentration

values were significantly higher than the values for simulated service conditions or at Ecorr in

0.1 M HCl solution. In addition, the hydrogen concentration increased with increasingly

negative charging potential for all the studied steels, such as for 980 DP, the hydrogen

concentration increased from 0.107 µg g-1

to 0.171 µg g-1

and up to 0.240 µg g-1

with the

charging over-potential changed from -0.257 V to -0.557 V and to -0.757 V. This trend was

consistent with the results from permeability experiments, as shown in Fig. 2.

Table 3 Diffusible hydrogen concentrations, Cd, in the studied steels under various hydrogen

conditions. Steel

designation

Conditions Charging potential

(VSCE)

Over-potential

(η, V)

Cd

(µg g-1

)

980 DP

uncharged - 0 0.002

at Ecorr in 3wt% NaCl - -0.039 0.031

in 0.1M NaOH -1.248 -0.257 0.107

in 0.1M NaOH -1.548 -0.557 0.171

in 0.1M NaOH -1.848 -0.857 0.240

1200 DP-GI

uncharged - 0 0.001

at Ecorr in 3wt% NaCl - -0.039 0.053

at EZn in 3wt% NaCl -0.998 -0.322 0.090

at Ecorr in 0.1M HCl - -0.203 0.070

in 0.1M NaOH -1.248 -0.257 0.120

in 0.1M NaOH -1.548 -0.557 0.218

in 0.1M NaOH -1.848 -0.857 0.300

980 QP

uncharged - 0 0.001

at Ecorr in 3wt% NaCl - -0.039 0.010

at EZn in 3wt% NaCl -0.998 -0.322 0.100

at Ecorr in 0.1M HCl - -0.203 0.085

in 0.1M NaOH -1.248 -0.257 0.197

in 0.1M NaOH -1.548 -0.557 0.252

in 0.1M NaOH -1.848 -0.857 0.347

Fig. 4 also provides a comparison between the measured diffusible hydrogen

concentration, Cd, and the prior measurements of the hydrogen concentration, Cd, from

permeability experiments, as presented in Fig. 2. Under simulated service conditions in 3%

NaCl solution, the diffusible hydrogen concentrations were similar to the values of total

hydrogen concentration determined by the permeability experiments, as was also the case for

that cathodically charged at over-potential of -0.257 V in NaOH solution. However, at more

negative charging over-potentials of -0.557 V and -0.857 V, the diffusible hydrogen

concentrations were lower than the total hydrogen concentration values obtained from our

permeability experiments.

9

Fig. 4 The measured diffusible hydrogen concentration, Cd, using the BRUKER G4

PHOENIX DH analyser in comparison with the hydrogen concentration, CH, determined

from permeability experiments.

4. Discussion

4.1. Hydrogen evolution

Under cathodic polarization or at the free corrosion potential, hydrogen evolution

reaction proceeds through reaction given by Eq. (1) or (2), and acts as part of the cathodic

partial reaction, balancing the anodic reaction. Ootsuka et al. [24] indicated that one mol of

hydrogen was absorbed for every 1000 mol of steel corroded, indicating that the hydrogen

evolution reaction compromised 0.1% of the cathodic reaction.

For the tests under cathodic polarization in 0.1 M NaOH and 3wt% NaCl solutions,

the electrons for hydrogen reduction reaction were provided by the external potentiostat,

whereas for those at free potential in 0.1 M HCl and 3wt% NaCl solutions, the electrons for

hydrogen evolution reaction came from the anodic reaction as following:

Fe → Fe2+

+ 2e (10)

Since the cathodic partial reaction as in Eq. (2) and the oxygen partial reduction

reaction balance the anodic partial reaction as in Eq. (10) in the 3wt% NaCl solution, a

change of either reaction rate causes an adjustment to minimize the effect of this change.

Cathodic polarization causes increased hydrogen evolution by Eq. (2) and decreased

corrosion rate by the rate of Eq. (10). This explains why the amount of corrosion for steels

cathodically polarized at EZn is significantly less than that at Ecorr in 3wt% NaCl solution, as

observed in our permeability experiments. During cathodic polarization in the 0.1 M NaOH

solution, hydrogen is produced due to the continuously provided electrons. A more negative

over-potential increased the current of the hydrogen evolution reaction as in Eq. (2), and

resulted in more hydrogen being produced.

In our study, for the experiments under cathodic charging in the 0.1 M NaOH solution,

prolonged charging was conducted. Thus, it was expected that the surface oxides were

reduced to a stable state, and there was equilibrium between the adsorbed hydrogen on the

hydrogen-entry side of the steel surface and the hydrogen dissolved into the steel [9]. For the

experiments in the 3wt% NaCl and 0.1 M HCl solutions, the same equilibrium was also

expected, since water molecule or hydronium ion diffused through gaps in the oxide film to

the steel surface and was reduced, releasing hydrogen adsorbed on the steel surface and then

introducing the adsorbed hydrogen into the specimen. Therefore, both of cathodic charging

10

and corrosion conditions in our study provided hydrogen evolution and the adsorbed

hydrogen on the steel surface was in equilibrium with the hydrogen dissolved into the steel,

so that hydrogen could be produced, be absorbed into the steel, permeate through and be

oxidized by the applied positive potential on the other side.

4.2. Permeation current density

For permeability experiments in the 3wt% NaCl and 0.1 M HCl solutions, the

adsorption of aggressive ions such as Cl- on the metal surface facilitates the corrosion rate of

the material and further the iron dissolution during corrosion depends on H+ ion concentration

more than on the Cl- ion concentration [25, 26]. Therefore, corrosion occurred in these

solutions at the free potential in this study and the corrosion rate in the 0.1 M HCl solution

was higher, reflected by the steeper initial slope of the permeation current density as in Fig.

2(b). As a consequence, more hydrogen was produced, entered into and diffused through the

specimen, leading to higher permeation rates and higher hydrogen concentrations at free

corrosion in 0.1 M HCl than in 3wt% NaCl, as shown in Fig. 2(b) and Table 3

For the experiments at EZn in the 3wt% NaCl solution, the cathodic polarization

increased the hydrogen reduction reaction, and decreased the corrosion of the steel, as stated

above. This explained the highest permeation rate and hydrogen concentration at EZn in 3wt%

NaCl solution, compared with the other conditions simulated auto service. During actual auto

service, the steel could be corroded, introducing hydrogen, and influencing the steel

properties, so the permeation experiments at Ecorr in 3% NaCl solution indicated how much

hydrogen had permeated under these conditions. In addition, in order to prevent the steel from

corrosion, steel can be galvanized, by applying a zinc coating on the steel surface. The

permeation experiments at EZn in the 3wt% NaCl solution indicated that this Zn coating could

significantly reduce the corrosion of the steel, but more hydrogen could be introduced into the

steel in addition to the hydrogen from the corrosion on the open surface of the steel. Thus,

more hydrogen would be introduced into and permeate through the galvanized steel if there is

access of the steel surface to the corroding solution.

As presented in Figs. 1 and 3, the permeation rate, i∞, for experiments under simulated

service conditions, comparing with cathodic charging conditions in NaOH solution, was

significantly lower, even much lower than those at -1.100 VHg/HgO, due to the lower rate of

hydrogen evolution reaction under simulated service conditions in NaCl solutions, despite

that the presence of Cl- that could accelerate the corrosion reactions. Besides, in the 0.1 M

NaOH solution, the permeation rate, i∞, increased with an increasingly negative potential, due

to the higher rate of hydrogen reduction reaction and a less potent trapping effect, which was

also indicated by the lower hydrogen trap site density as in our prior study [11]. This trend is

consistent with that in other research [9, 22].

4.3. Service versus cathodic charging

Fig. 2 presents the relationship between the total hydrogen concentration at the input

side of the steel specimen, CH, and the over-potential, η, for all the steels under simulated

service conditions, and under cathodic charging in the 0.1 M NaOH solution, determined

through permeability experiments. This allows a direct comparison between the hydrogen

concentrations produced under service conditions with those under cathodic charging

conditions. A turning point at over-potential of around -0.5 V was observed in the

relationship between CH and η for each steel under cathodic charging, as was also reported by

Bockris et al. [27] and Liu et al. [9]. This turning point was attributed to the change of

hydrogen evolution mechanisms from coupled discharge-recombination mechanism at lower

over-potentials to a slow discharge-fast electrochemical mechanism at higher over-potentials,

meaning that the electrochemical discharge of adsorbed hydrogen became more likely.

11

Fig. 2 shows that, for each steel, the hydrogen concentration at the input side of the

steel, CH, was significantly lower under simulated auto service conditions in 3wt% NaCl or in

0.1 M HCl solution than that under cathodic charging in 0.1M NaOH solution, even lower

than that at the least negative charging potential of -1.100 VHg/HgO, as was also the case for

hydrogen permeation current density. And further, extrapolating the experimental data

obtained in the 0.1 M NaOH solution, the CH value for all the steels in 0.1 M NaOH solution,

at each over-potential corresponding to that under simulated service corrosion conditions, was

also higher than that experimentally determined under simulated auto service conditions. This

indicates that at the same over-potential, more hydrogen entered in and permeated through the

steel specimen in the 0.1 M NaOH solution than that in the 3wt% NaCl solution. This is

considered attributed to the lower rate of hydrogen evolution reaction under simulated service

conditions in NaCl solutions, as also stated above in section 4.2.

4.4. Diffusible and total hydrogen content

Fig. 4 and Table 3 show that the amount of diffusible hydrogen in the as-received

steels was quite low, indicating that either there was little hydrogen in the steels due to steel

production, or that the hydrogen had effused out of the steels by the time that these steel

samples were tested in this work.

Fig. 4 also compares the amount of diffusible hydrogen, Cd, measured by hot

extraction analyser, with the total amount of hydrogen, CH, as measured by the permeability

experiments. There was good agreement between Cd and CH for the conditions of (i) free

corrosion in 3wt% NaCl solution, (ii) free corrosion in 0.1 M HCl, (iii) cathodic polarisation

of the steel to the zinc potential in 3wt% NaCl, and (iv) cathodic polarisation at the lowest

over-potential value of -1.100 VHg/HgO in the 0.1M NaOH solution. This good agreement

provides validation of using Eqs. (7) to (9) to analyse the permeation curves and estimate the

total hydrogen concentration, CH. Besides, Fig. 4 shows that, at the larger values of over-

potential, such as -0.557 V and -0.857 V, for cathodic hydrogen charging, the measured

amount of diffusible hydrogen, Cd, was somewhat less than the total amount of hydrogen, CH,

attributed to the loss of some diffusible hydrogen during the time between the end of the

cathodic charging and the start of measuring the amount of diffusible hydrogen. This delay

was inevitable since the specimen needed to be washed, dried, inserted into the measurement

apparatus and the commencement of heating for hydrogen extraction. The loss of diffusible

hydrogen by diffusing out of the sample is related to the amount of diffusible hydrogen and

was thus much larger at the higher hydrogen concentrations. In addition, as shown in prior

study [11], the effective diffusion coefficient tends to be larger at the higher concentrations

because of less trapping [9, 22], and consequently the diffusion loss is greater at the higher

hydrogen concentrations. In contrast, the total amount of hydrogen, CH, was measured by the

permeation technique under steady state conditions, and thus represented the total hydrogen

in the specimen.

5. Conclusions 1. The values of the effective diffusion coefficient measured in our work were consistent

with literature values.

2. The hydrogen permeation rates and hydrogen concentration increased with increasingly

negative charging potential in 0.1 M NaOH solution.

3. The hydrogen permeation rates and hydrogen concentrations for simulated auto service

conditions were significantly lower than the values at the least negative cathodic charging

potential of –1.100 VHg/HgO in 0.1 M NaOH solution.

4. The Zn coating on the steels could protect the substrate steel from corrosion, however,

once the Zn coating was corroded under actual service condition, more hydrogen could

enter and permeate through the steel.

12

ACKNOWLEDGEMENTS

This research is supported by the Baosteel-Australia Joint Research & Development Centre

(BAJC) Grant, BA13037, with linkage to Baoshan Iron and Steel Co., Ltd of China.

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