ORIGINAL RESEARCH ARTICLE

High-Temperature Hydrogen Attack on 2.25Cr-1Mo
Steel: The Roles of Residual Carbon, Initial
Microstructure and Carbide Stability

M.A.M. ALSHAHRANI, S.W. OOI, M. HÖRNQVIST COLLIANDER,
G.M.A.M. EL-FALLAH, and H.K.D.H. BHADESHIA

High temperature hydrogen attack is a damage mechanism that occurs in critical steel
components in petrochemical plants and refineries when the hydrogen penetrates the steel and
reacts with the carbides within to produce pores containing methane. With the motivation of
understanding the role of carbide stability on the reaction with hydrogen, samples of a classic
214Cr-1Mo steel were subjected to a variety of heat treatments that generate a corresponding
variety of precipitates, prior to exposure to high-pressure hydrogen in an autoclave. Using
quantitative carbide, porosity and microstructural characterisation, it has been possible to
demonstrate the roles of four variables: (a) the carbon residue present in the ferrite; (b) the
non-equilibrium chemical composition of carbide; (c) the fraction of the carbide that is closest to
the thermodynamic equilibrium state and (d) the initial microstructural state, i.e., whether it is
martensitic or bainitic prior to heat treatment.

https://doi.org/10.1007/s11661-022-06809-9
� The Author(s) 2022, corrected publication 2023

I. INTRODUCTION

HYDROGEN embrilttlement is one of the mecha-
nisms that threatens the integrity of steel equipment in
the oil and gas industry. The atomic hydrogen is the
cause of the embrittlement.[1] The dissociation of molec-
ular hydrogen into nascent atomic hydrogen at the steel
surface commonly occurs by the chemical reaction
caused by corrosion or cathodic protection at the steel
surface. The nascent hydrogen atoms fail to recombine
at the surface due to the presence of poisoners such as
hydrogen sulfide and cyanides. The atomic hydrogen

then diffuses into the steel.[2,3] However, in the absence
of an aqueous environment, i.e., high-temperature
environments, the corrosion reaction is eliminated at
the surface. Molecular hydrogen then dissociates under
the high pressure of the environment at the steel surface,
which is the hydrogen dissociation mechanism in
high-temperature, high-pressure hydrogen attack that
is discussed in this paper.
Apart from accumulating at stress-fields such as

inclusions and ferrite/carbide interfaces, the diffusable
atomic hydrogen reacts chemically with the carbon that
is tied up in carbides to form methane bubbles which
become a permanent feature of the steel as they are too
large to diffuse out, as seen in the following reaction:
C+2H2 $ CH4. The methane reaction occurs prefer-
entially at the least stable carbide, cementite, as follows:
Fe3C+2H2 $ CH4 +3Fe. The reaction can still take
place to a lesser extent at other substitutionally alloyed
carbides that are precipitated in creep-resisting steels,
for example, M2C, M7C3, M23C6 etc., where ‘M’ stands
for a mixture of metal atoms. The formation of
methane-containing bubbles leads to internal stresses
that weaken grain boundaries, which eventually lead to
the formation of microcracks and ultimate failure under
the influence of the external stresses that the component
is designed to support or even by the local creep
introduced by the bubble formation.[4–16] It could be
argued, therefore, carbides that are more thermody-
namically stable in steel should result in an enhanced
resistance to hydrogen attack.[17,18] The problem is

M.A.M. ALSHAHRANI is with the Materials Science and
Metallurgy, University of Cambridge, Cambridge CB2 3QZ, UK
and also with the Saudi Aramco Research and Development Centre,
Dhahran 34466, Saudi Arabia. Contact e-mail:
mohammed.shahrani.34@aramco.com S.W. OOI is with the
Materials Science and Metallurgy, University of Cambridge. M.
HÖORNQVIST COLLIANDER is with the Department of Physics,
Chalmers University of Technology, Göteborg 41296, Sweden.
G.M.A.M. EL-FALLAH is with the School of Engineering,
University of Leicester, Leicester LE1 7RH, UK. H.K.D.H.
BHADESHIA is with the Materials Science and Metallurgy,
University of Cambridge and also with the School of Engineering
and Materials Science, Queen Mary University of London, London E1
4NS, UK.
Manuscript submitted May 13, 2022; accepted August 21, 2022.
Article published online October 8, 2022

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 53A, DECEMBER 2022—4221

http://crossmark.crossref.org/dialog/?doi=10.1007/s11661-022-06809-9&domain=pdf


important because failures of the steel in petrochemical
plant can on rare occasions lead to fatal catastrophes
such as at the Tesoro Anacortes Refinery;[19] more
commonly, the attack requires the costly replacement of
components.[20]

The current best practice for design against hydrogen
attack uses an experience-based approach, expressed as
the Nelson curves,[21] Figure 1.[22] These curves define the
temperature and hydrogen pressure at which established
steels can operate without the risk of hydrogen attack.
The curves suggest that carbides-based Cr, Mo and V
are less likely to be reduced by hydrogen.[23,24] However,
although the average composition is clearly an impor-
tant factor, it is possible in each of these steels to
generate a variety of microstructures so the Nelson
curves do not represent the full picture. Heat treatments
could play a major role in precipitating specific carbides
that resist hydrogen attack. There is a general feeling
that design based on a better understanding of the
mechanism and kinetics of attack may lead to better
utilisation of existing steels or the development of new
ones that better resist the attack.[14,25]

Studies on a commercially used 214Cr-1Mo steel have
shown that the equilibrium carbide, M23C6 (‘M’ stands
for metal atoms) is resistant to hydrogen attack when
compared with cementite; however, there are discrepan-
cies in the perceived order of stability in mixtures of
M7C3, M2C and M6C.

[4,26]

The present work assesses the 214Cr-1Mo steel in a
variety of microstructural states in order to explicitly
reveal the hydrogen-reaction resistance of precipitated
carbides. Quantitative methods based on synchrotron
X-ray diffraction and microscopy have been used to
characterise samples exposed to high temperature,
high-pressure hydrogen. It will be shown that the
experimental data can be consistently interpreted within
a thermodynamic framework of phase stabilities.

II. EXPERIMENTAL PROCEDURE

A. Material and Heat Treatment

The steel was supplied by British Petroleum. It was
electroslag-remelted and then forged to produce a
homogeneous component. There was, therefore, nothing
in the metallography of the samples that indicated
chemical segregation. All the samples were cut from
adjacent positions so we do not expect variations in
inclusion density between the samples.
Carbon is usually present within the steel in the form

of carbides. The work here is focused on relating the
depletion of carbides due to hydrogen attack on
214Cr-1Mo steel to the heat treatment, tempering condi-
tions and deviation from equilibrium. The composition
of the actual steel used is listed in Table I.
The heat treatments were selected based on the work

of Baker and Nutting.[27,28] A number of combinations
of carbides could be generated to reveal their influence
on hydrogen attack resistance. Five specimens, cubes
with a side length of � 2:5 cm, were austenitised for 2 h
at 960 �C. Two of the cubes were then quenched into
water with the remainder air cooled. Each cube then was
subjected to a different heat-treatment in order to induce
a variety of carbide mixtures; we note that the carbide
compositions in terms of the metal atoms are also
sensitive to the heat-treatment temperature. The details
are listed in Table II.

B. Hydrogen Exposure

Judging from Figure 1, the 214Cr-1Mo steel is expected
to be susceptible to attack if the temperature exceeds
450�C when the hydrogen partial pressure exceeds
13.7 MPa. The equipment used was at The Welding
Institute (TWI), which can achieve up to 525 �C and a
hydrogen pressure of 10 MPa, the precise conditions
used in the present work. The TWI autoclave was used

Fig. 1—The Nelson curves used for designing against hydrogen attack. The area below each curve represents the safe operating zone for the
corresponding steel. In red is the 2.25Cr-1Mo steel studied in this work. Adapted from Ref. [22] (Color figure online).

4222—VOLUME 53A, DECEMBER 2022 METALLURGICAL AND MATERIALS TRANSACTIONS A



at maximum capacity for 20 days to be consistent with
the work of Wanagel et al.,[29] although the parameters
possible on the TWI autoclave are less severe that those
in Reference 29. It also is costly to use the facility so
longer exposures could not be accommodated. The
specimens were polished before hydrogen exposure with
no visible signs of oxide, and then placed into the
hydrogen chamber. There were no differences found in
the surfaces after exposure, between the different heat
treatments.

C. Characterisation

The characterisation focused on the comparison of
results before and after hydrogen exposure to give a
clear picture of damage evolution:

� as the carbide reacts with hydrogen, it leaves behind
a methane-filled void which can be imaged and
measured using microscopy on unetched specimens.
The images were then processed using ImageJ
software to determine the area fraction occupied
by voids.
The FEI Tecnai Osiris FEG TEM, fitted with an
energy dispersive spectroscopy detector, was used
for microstructural characterisation of fine carbides,
as well as carbide identification. The microscope has
an accelerating voltage of 200 kV. Carbide reflec-
tions were identified using diffraction patterns, while
their morphology and orientation were revealed with
dark-field imaging. Scanning transmission micro-
scopy was used for high resolution imaging of fine
carbides.

� Synchrotron X-ray analysis was used to measure
how the carbide volume fractions have changed
following hydrogen exposure. The diffraction inten-
sities using synchrotron X-ray can be sufficiently
large to measure the relatively small fractions of
carbide present in each of the specimens.
Experiments were conducted on The Swedish Mate-
rials Science Beamline P21.2 at PETRA III in the
Deutsches Elektronen-Synchrotron. The X-ray

wavelength was 0.2066 Å, monochromated using
diffraction. X-ray detection was done using a
PILATUS3 X CdTe 2M detector with pixel size of
172 lm� 172 lm. The sample detector distance was
determined to be 2800 mm to allow the largest
possible diffraction at 2h<15�, where the most
intense peaks of carbides are expected. The X-ray
beam size was chosen to be 0.5 mm 9 0.5 mm with
a dwell time of 5 s per measurement point. Hundreds
of images were captured, spaced out evenly across
each specimen. Analysis was performed using
Python modules; pyFAI-calib2 for detector calibra-
tion and pyFAI-integrate for data reduction. Riet-
veld analysis using MAUD[30] was used to calculate
the volume fraction of carbides.

� Energy dispersive X-ray analysis in a transmission
electron microscope was used to identify the carbides
and measure their compositions.

III. RESULTS AND DISCUSSION

A. Carbides Precipitated Before Hydrogen Exposure

The results of synchrotron X-ray analysis are pre-
sented in Figure 2.
Figure 3 compares the Baker and Nutting findings on

the types of carbide observed, with the data determined
using synchrotron X-ray analysis for the specific steel
used here. Although there is broad agreement, only
Sample A completely complies with the original work. It
is speculated that discrepancies for other specimens
might be because of the detailed difference in the heat
treatment, for example the cooling rate during quench-
ing or air cooling, which is affected by the size of the
specimen. Nevertheless, the purpose of using different
heat treatments was to precipitate a variety of combi-
nations of carbides and that was achieved.
Phase stability calculations using ThermoCalc

(TCFE10 database)[31] were used to estimate equilibrium
carbide volume fractions and their chemical composi-
tions. It is appreciated that these represent

Table I. Chemical Composition (Weight Pct) of the 21
4
Cr-1Mo Used

C Si Cr Mo Ni Mn P S Cu Al V

0.14 0.1 2.24 0.97 0.37 0.42 0.007 0.001 0.04 0.043 0.005

Table II. Heat Treatments Performed on Different 21
4
Cr-1Mo Specimens

Sample Heat Treatment Carbides Expected[28]

A Quench, temper @ 650�C, 50 h M6C + M2C + M7C3

B Quench, temper @ 700�C, 1 h M3C + M2C+ M7C3

C Air-cool, temper @ 650�C, 50 h M6C + M2C + M7C3

D Air-cool, temper @ 700�C, 50 h M23C6 + M6C + M2C + M7C3

E Air-cool, temper @ 750�C, 5 h M23C6 + M6C + M2C + M7C3

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 53A, DECEMBER 2022—4223



thermodynamic equilibrium rather than kinetics, but as
will be seen later, they nevertheless can give useful
information in interpreting complex observations.
Figure 4 shows that for all the tempering temperatures
investigated in the present work, the only carbide to be
able to exist in equilibrium with ferrite is M23C6. All
other carbides that are observed in the steel therefore
must be regarded as transient. It still is useful to
characterise the equilibrium for these transient carbides.

This was done in the calculations by suppressing the
most stable carbide in the precipitation sequence to
allow the transients to appear; these calculations are
henceforth referred to as constrained equilibria.
Table III compares the carbide volume fractions

expected by ThermoCalc and those measured using
synchrotron X-rays before hydrogen exposure. The
comparison shows that with the heat treatments imple-
mented, most of the carbides did not reach the expected

Fig. 2—Synchrotron X-ray line profiles. (a) Sample A. Note that the scales for all the remaining profiles are identical to those of sample A. (b)
Sample B. (c) Sample C. (d) Sample D. (e) Sample E.

4224—VOLUME 53A, DECEMBER 2022 METALLURGICAL AND MATERIALS TRANSACTIONS A



equilibrium nor constrained-equilibrium volume frac-
tions. The exceptions being samples D and E in which
the M23C6 and M6C both are approximately at equilib-
rium. This is because those samples were tempered at the
highest temperatures. Table IV shows the equilibrium
composition of carbides at the tempering temperatures
used in this work.

As emphasised earlier, only the equilibrium M23C6

values were obtained while allowing all the other
carbides to exist in the calculations. The same applies
to the data for Sample D, where the volume fraction of
M6C is very small and the agreement between the
thermodynamically calculated and actual fraction of
M23C6 is consequently good. On the other hand, M7C3

and M6C values could only be estimated after suppress-
ing M23C6. M2C and cementite values were obtained
similarly after suppressing M23C6 and M7C3.

Therefore, the comparison between the equilibrium
thermodynamic estimates and measured data are strictly
only valid for Sample E which contained just M23C6; the
agreement there is excellent. The rest of the calculations
represent constrained equilibria, which if given sufficient
time at temperatures would eventually reach equilib-
rium, but have not done so with the heat treatments T

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Fig. 4—Equilibrium calculations for the alloy listed in Table I. M3C,
M6C, M23C6, M2C, M7C3, where ‘M’ represents the metal atoms
Fe, Cr, Mo, were all allowed to exist in the calculations, together
with the residual phase a. No carbide other than M23C6 was
predicted at equilibrium.

Fig. 3—Carbides as expected from Nutting[28] and as found here. In
red are the current findings with samples represented by their
designated letters (Table V) (Color figure online).

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 53A, DECEMBER 2022—4225



implemented. In contrast, Samples A-C show
metastable microstructures so closure is not expected
with the equilibrium calculations.

B. Void Fractions After Hydrogen Exposure

In the initial state, the steel contains only non-metallic
inclusions but no voids. The inclusion fraction was
measured prior to hydrogen exposure, using microscopy
imagingofunetched specimens thenanalysedusing ImageJ
software, as VV ¼ 8 � 10�5 � 9 � 10�6. The same
method was used to characterise the increase in the void
volume fraction following hydrogen attack. The increase
was then identified as voids introduced during hydrogen
exposure. A summary of the results is listed in Table V.
Sample C has the highest VV, which suggests that it
underwent the highest hydrogen attack rate with an
increase of two orders of magnitude after hydrogen
exposure. This was followed by sample D, A+B in terms
of descending VV, with the least damage in sample E.

Samples A, B, D and E had an increase inVV by one order
of magnitude in comparison to the state prior to exposure.
The increase in void volume fraction following

hydrogen exposure (Table V) might correlate with the
reduction of carbide volume fractions. To investigate
how the carbides resisted hydrogen attack, synchrotron
X-ray analysis was carried out on the hydrogen-exposed
specimens to permit the accurate measurement of small
volume fractions.

C. Carbide Fractions Following Hydrogen Exposure

A representation of how the synchrotron X-ray
analysis was carried out on the specimens is shown in
Figure 5. Measurements were conducted horizontally
from surface to surface. After the whole sample has been
examined, all values from the same position were
averaged to give one spectrum for that location, which
was then analysed using Rietveld analysis to estimate
carbide fractions at that position. The first and last

Table IV. ThermoCalc Prediction of Equilibrium Chemical Composition (Weight Pct) for Carbides as a Function of Temperature

for 21
4
Cr-1Mo Steel

M3C M2C M6C

650 �C 700 �C 750 �C 650 �C 700 �C 750 �C 650 �C 700 �C 750 �C

C 6.86 6.83 6.81 6.56 6.79 0.02 2.52 2.56 2.59
Cr 40.70 35.52 30.97 11.57 15.09 2.74 2.93 3.44 3.76
Mo 1.04 1.22 1.44 79.70 74.51 0.79 66.69 64.17 61.69
Fe 48.38 54.01 58.94 1.83 3.36 94.68 27.38 29.32 31.45
Mn 2.98 2.36 1.78 0.32 0.23 0.89 0.00 0.00 0.00

M7C3 M23C6

650 �C 700 �C 750 �C 650 �C 700 �C 750 �C

C 8.41 8.41 8.40 5.05 5.07 5.10
Cr 54.32 50.90 47.17 37.87 34.31 30.69
Mo 10.50 9.80 9.38 18.44 16.94 14.69
Fe 24.33 28.79 33.32 38.14 43.12 48.91
Mn 2.41 2.07 1.68 0.33 0.38 0.43

The composition of M23C6 is when all carbides were permitted to exist but only M23C6 was predicted in association with a. Those of M7C3 and
M6C are for the case where M23C6 was suppressed. The compositions of M2C and cementite represent the case where M2C, cementite and M6C were
only allowed to exist.

Table V. Void or Inclusion Volume Fractions Before and After Hydrogen Exposure

Unexposed Sample
Inclusion Volume Fraction
8� 10�5 � 9� 10�6

Sample Void+inclusion volume fraction
A (Quenched, ! 650 �C for 50 h) 47� 10�5 � 2� 10�5

B (Quenched, ! 700 �C for 1 h) 46� 10�5 � 2� 10�5

C (Air cooled, ! 650 �C for 50 h) 120� 10�5 � 3� 10�5

D (Air cooled, ! 700 �C for 50 h) 52� 10�5 � 2� 10�5

E (Air cooled, ! 750 �C for 5 h) 33� 10�5 � 3� 10�5

The quantities for the exposed specimens include the inclusion fraction from the unexposed state. The error bars were calculated by dividing the
standard deviation by the square root of the number of measurements.

4226—VOLUME 53A, DECEMBER 2022 METALLURGICAL AND MATERIALS TRANSACTIONS A



positions are also the nearest with respect to the
hydrogen-exposed surface.

According to the data in Table V, Samples A and B
were the 2nd least damaged specimens in terms of the
reduction in carbide volume fractions. Synchrotron X-ray
analysis showed no noticeable reduction in carbides
except for one position for M7C3 in sample B. M2C and
M6C showed good resistance to reaction with hydrogen
as their fractions remained stable, explaining the good
overall-resistance of those two specimens.

Sample C exhibited the least resistance to hydrogen
attack. The void volume fraction increased by two
orders of magnitude relative to the unexposed specimen,
Table V. Figure 6 shows that the reduction was greatest
for M23C6 in sample C with its volume fraction reduced
from 0.0113 ± 0.0008 in the unexposed specimen to as
low as 0.0051 ± 0.0005 in one of the positions in the
exposed specimen, a reduction of 0.0062. Figure 6 shows
that M7C3 also decreased from 0.0048 ± 0.0007 in the
unexposed specimen to 0.0021 ±0.0006 at one of the
positions in the exposed specimen. M2C and M6C
fractions remained essentially unchanged.

M6C also resisted attack in sample D. The increase in
VV in this specimen is caused by the reduction of M23C6,
which was reduced by 0.0056, a similar reduction to that
observed in sample C., as seen in Figure 7. This
specimen was ranked 2nd in terms of susceptibility to
hydrogen attack.
The smallest increase in VV was observed in sample E,

which contained only M23C6. The high tempering
temperature (750 �C) influences the composition of
M23C6 making it closer to its equilibrium composition,
Table VI, which will be discussed later. This explains the
excellent hydrogen attack resistance in this specimen.
It is interesting to note that a greater degree of

carbide-depletion is observed towards one of the spec-
imen surfaces in Figures 6 and 7. The specimen posi-
tioning, Figure 5, suggests that the access to constant
high hydrogen pressures is less to the bottom of the
specimen than other areas, which could introduce a
difference in carbide-depletion amounts. Those heavily
depleted regions correspond to the top end of the
specimens in Figure 5, where access to hydrogen pres-
sure is continuously greater.
The voids volume fraction can be seen to correspond

to the depletion of carbides by Figure 8 that shows the
gradient of voids area percentage across specimen C.

Fig. 5—A representation of how measurements are taken with
synchrotron X-ray for hydrogen-exposed specimens. The sketch on
the right shows how the square-sectioned sample was placed within
the autoclave during hydrogen exposure. As a result, position 1 has
less access to hydrogen than position n.

Fig. 6—Sample C. The fractions prior to exposure are represented by the blue lines. First and last positions represent the closest proximity to
the exposed surface. M2C and M6C fractions remained essentially unchanged, so are not plotted here (Color figure online).

Fig. 7—Sample D. The fraction prior to exposure is represented by
the blue line. First and last positions represent the closest proximity
to the exposed surface (Color figure online).

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 53A, DECEMBER 2022—4227



Each square represents a micrograph that was analysed
by ImageJ to calculate the voids volume fraction. Note
that the top of Figure 8 corresponds to the top of
specimen when it was positioned in the autoclave as
shown in Figure 5. It can be seen from Figure 8 that
more voids have formed in the vicinity of the edges of
the specimen, where exposure to hydrogen is expected to
be the highest. It can also be seen that voids have formed
at the top part of the specimen more than its bottom,
where access to hydrogen pressure is continuously
greater due to specimen positioning showed in
Figure 5. The gradient of voids being the most towards
the edges of the specimen shows the correlation between
the voids formation and hydrogen ingress into the steel.

In summary, Mo-rich carbides, M2C and M6C, seem
to resist attack under all the conditions studied more
than Cr-rich carbides, where resistance was more
dependent on the extent of tempering. M7C3 and
M23C6 resistance varied between samples. Therefore,
hydrogen attack resistance cannot be decided on the
basis of the carbide type alone. It may depend on other
factors such as heat treatment, i.e., quenched or
normalised, and tempering temperature and duration.
M7C3 in 214Cr-1Mo steel are transient carbides

(Figure 4) and to eventually be replaced by M23C6,
which explains why its fraction decreases. In order to
investigate Cr-carbides further, the behaviour of M23C6

was examined as it showed low resistance in samples C
and D, but high resistance in sample E.

D. Effect of Deviation from Equilibrium

M23C6 showed poor hydrogen attack resistance in
samples C and D as it depleted by about 0.006 of its
original volume fractions, while the hydrogen attack
resistance of M23C6 was excellent in sample E. Sam-
ples C and D were tempered for 50 h at 650 �C and
700 �C, respectively, while sample E was tempered at
750 �C for 5 h. Therefore, the hydrogen attack resis-
tance of M23C6 increased as the tempering temperature
increased. In terms of equilibrium, calculations showed
that M23C6 should be the only carbide present at
equilibrium at all tempering temperatures (Figure 4).
Not surprisingly, this suggests that only sample E was at
equilibrium as it contained only M23C6.
For further investigation, the chemical composition of

M23C6 was calculated at equilibrium using ThermoCalc
and measured experimentally using energy dispersive
X-ray analysis in a transmission electron microscope.
The results are listed in Table VI; the chromium
concentration within M23C6 is clearly much less than
equilibrium for samples C and D. Parathasarathy and
Shewmon investigated the correlation between hydrogen
attack rates and carbon activity against the atom
fraction of chromium in the matrix ferrite,[32] in which
they concluded that a high chromium content in the
matrix indicates a low iron content, hence a reduced
carbon activity, which must lead to a greater hydrogen
attack resistance, Figure 9. The calculations are based
on the estimation of the carbon activity that is in
equilibrium with M23C6 using the reaction

yCrCrþ
6

23
Cþ yFeFe ¼ ðFeyFeCryCrÞC6=23

where yFe ¼ xFe=ðxFe þ xCrÞ, with x representing the
mole fraction of the solute identified in the subscript.
It follows that yCr ¼ 1� yFe. The standard-free energy
change for this reaction was compiled from thermody-
namic data[33,34] as

Table VI. Equilibrium and Measured Chemical Composition (Weight Pct) of M23C6 at Different Tempering Temperatures

Sample C, air cool ! 650�C, 50 h Sample D, air cool ! 700�C, 50 h Sample E, air cool ! 750�C, 5 h

Calculated EDX Calculated EDX Calculated EDX

C 5.052 5.071 5.107
Cr 37.877 15.61 ± 0.20 34.317 16.01 ± 0.13 30.691 27.43 ± 0.05
Mo 18.448 6.62 ± 0.13 16.943 3.93 ± 0.07 14.692 4.90 ± 0.03
Fe 38.146 70.04 ± 0.42 43.129 71.79 ± 0.26 48.916 58.33 ± 0.07
Mn 0.332 2.68 ± 0.10 0.385 3.19 ± 0.05 0.435 4.23 ± 0.02

Carbon content used was calculated by ThermoCalc and EDX measurements for other elements were corrected accordingly.

Fig. 8—A representation of the cross section of specimen C that
underwent the highest attack. Each square represents a micrograph
that was analysed by ImageJ. Higher voids formation in the vicinity
of the edges is assumed to correlate to the hydrogen ingress into the
specimen. Voids have formed more towards the top of the specimen
due to the greater access to hydrogen pressure because of the
positioning showed in Fig. 5.

4228—VOLUME 53A, DECEMBER 2022 METALLURGICAL AND MATERIALS TRANSACTIONS A



DG� ¼yCr

�
G�

CrC6=23
þ G�

Cr �
6

23
G�

C

�

þ yFe

�
G�

FeC6=23
þ G�

aFe �
6

23
G�

C

�

þ RTðyFe lnfyFeg þ yCr lnfyCrgÞ
þ yFeyCr1100 Jmol�1

so that the equilibrium constant K in terms of the
activities ai is given by

K ¼ exp

�
�DG�

RT

�
¼

aMC6=23

a
yCr
Cr � a

yFe
Fe � a

6=23
C

:

Using this, Parathsarathy and Shewmon went on to
isolate the activity of carbon as

aC ¼ exp

�
DG�

RT

�
aMC6=23

ðayCrCr � a
yFe
Fe Þ

23=6 ½1�

but the equation is incorrect; it should in fact be

aC ¼
"
exp

�
DG�

RT

�
aMC6=23

a
yCr
Cr � a

yFe
Fe

#23=6

: ½2�

It has been verified that the original error was typo-
graphical, their results are correct as indicated by our
calculations in Figure 9(a) which compares well with
Figure 5 of Reference 32, showing how the activity of
carbon in the matrix varies with the matrix concentra-
tion of chromium and the ratio yFe=yCr. Since the
temperature they used for their calculations differed
slightly from our experiments, Figure 9(b) shows the
results of relevant behaviour for 525 �C, using
Equation [2].

The data in Table VI show that the Fe and Cr
concentrations are not close to equilibrium when tem-
pering at 650 �C to 700 �C, because the measured Cr
concentrations should be much larger, making the
M23C6 more reactive at the lower temperatures. This

explains the difference in hydrogen attack behaviour of
M23C6 when generated using different tempering tem-
peratures. Therefore, if equilibrium conditions for
M23C6 are reached for low tempering temperatures
(650 �C), the carbon activity would be minimised,
thereby decreasing the tendency to react with hydrogen
and form methane. But unfortunately, the tempering
times required would be impractical as literature has
shown that not even 500 h at 700 �C would ensure
equilibrium conditions,[32] unless the temperature is
increased as seen in sample E (750 �C).
It is interesting to mention that equilibrium Cr

concentrations decrease with increasing tempering tem-
perature, as seen in Table VI, suggesting that greater
concentrations of Cr dissolved in ferrite that is tempered
at a higher temperature.
Table VI also shows that the molybdenum concen-

tration in M23C6 is much smaller than consistent with
equilibrium, for all the samples listed. However, this
measured concentration of molybdenum is similar for
all the samples listed, so probably does not play are role
in differentiating the hydrogen attack behaviour of the
three samples.

E. Effect of Microstructure

Samples A and C were both tempered at 650 �C for
50 h. Sample A was quenched to martensite before
tempering while sample C was air cooled into a bainitic
state before tempering. Sample C has greater hydrogen
attack damage than sample A (Table V). This difference
cannot be attributed to the tempering treatment, but
possibly to microstructure prior to tempering. Of
course, sample E also had a bainitic initial microstruc-
ture but it was tempered much more severely to almost
reach equilibrium, so its excellent attack resistance is
clearly a consequence of the high chromium concentra-
tion in its sole carbide, M23C6.
The carbides expected in both specimens C, D are

M7C3, M6C and M2C, but also M23C6 in sample D
(Figure 3). However, synchrotron X-ray indicated that

(a) (b)

Fig. 9—Calculations using Equation 2. (a) The relationship of carbon activity with the atom fraction chromium in the matrix for various atom
fractions of iron in the carbide metal content in 214Cr-1Mo steel for the temperature used by Parthasarathy and Shewmon.[32] Adapted from Ref.
[32] by permission from Springer Nature. (b) Calculations for the temperature representative of the present experiments.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 53A, DECEMBER 2022—4229



M23C6 was precipitated in sample C along with the
expected carbides, while sample A contained only the
transition carbides.

To understand why alloy carbide precipitation is
slower in the initially martensitic microstructure, it is
important to understand that the bainite- and marten-
site-start temperatures of the steel are 547 and 421 �C,
respectively, so the bainite forms at a temperature where
there is significantly greater atomic mobility. Tsai and
Yang[35] conducted detailed electron microscopy for a
sample of 2 1

4Cr-1Mo steel tempered at 700 �C to show
that alloy carbides form more rapidly when the initial
state is bainitic. This might be expected from the higher
transformation temperature, the microstructure in effect
being partly tempered during transformation. Further-
more, the martensite is expected to contain a greater
density of dislocations; the formation of cementite
during the early stages of tempering would therefore
be retarded because carbon is more stable at dislocations
than in cementite.[36] Given that the volume fractions of
the transient alloy carbides as measured in sample A are
almost identical to the corresponding carbides in sam-
ple C, it must be concluded that sample A has not
completed the precipitation process and this applies also
to sample B.

If the total amount of carbon that is in the form of
carbides (wcarbide), is calculated, then using the experi-
mentally measured data from Table III and assuming
stoichiometry,

wcarbide =wt pct wresidue =wt pct abefore = Å aafter = Å

Sample A 0:106� 0:007 0:038� 0:007 2:868 2:869

Sample B 0:098� 0:007 0:042� 0:007 2:869 2:869

Sample C 0:121� 0:013 0:019� 0:013 2:868 2:868

Sample D 0:146� 0:012 0:0 2:868 2:868

Sample E 0:140� 0:011 0:0 2:868 2:868

½�

where wresidue represents the difference between the
average carbon concentration (0.14 wt pct ) and that
tied up in carbides. The residual concentrations might
seem large given the high-temperature tempering expe-
rienced by all of these samples. However, in a
0.17C-9.3Cr-1.6Mo-0.27V wt pct steel cooled to marten-
site and then tempered at 590 �C for 8 h and then at
700 �C for 16 h, atom probe measurements revealed the
tempered ferrite to contain 0.035 wt pct of carbon
within the ferrite matrix that results from the tempering
of martensite.[37] After creep at 600 �C for 14522 h at
160 MPa, the carbon concentration in the ferrite
decreased to 0.0151 wt pct,[38] which still is much greater
than consistent with the equilibrium of alloy carbides
with ferrite.

It is likely that any excess carbon is located at defects
rather than in solid solution; the lattice parameters do
not show significant variation between the samples in
both the initial (abefore) and exposed (aafter) exposed
states; the error bars in the measurements of lattice
parameters are less than � 0:001 Å.

If the carbon is located at defects such as dislocations,
then the martensite should contain more excess carbon
than bainite, because of its greater dislocation density,
explaining the larger residual concentration in samples
A and B.
It is clear from the samples studied that both the

residual carbon concentration in the ferrite, and the
fraction of M23C6 precipitation play a role. The multiple
regression analysis illustrated in Figure 10 shows that
these two variables can explain a large amount of the
observed variation in the tendency to form voids, albeit
with a correlation coefficient of only 0.74. If Vtotal

V and

Vvoid
V represent the inclusion+void and void volume

fractions, respectively, then the corresponding regres-
sion equations are:

Vtotal
V

10�5
¼481� 16266VM23C6

V � 10695wresidue � 33

and Vvoid
V ¼Vtotal

V � 8� 10�5

where wresidue is in weight percent and the relationships
applicable over the ranges studied here.
Probably, the most important missing variable is the

chromium content of the M23C6, which clearly will
influence the activity of the carbon in the matrix,
Figure 9. Since this carbide did not form in samples A
and B, there are only three data points available for
samples C-E in Table VI. Since multiple regression with
three variables and three data is not possible, the Vtotal

V
was regressed against the chromium concentration of
M23C6,

Vtotal
V

10�5
� �4:7wM23

wt pct CrC6 þ 161 ½3�

with a correlation coefficient of 0.7. The negative
coefficient is consistent with a greater chromium con-
centration reducing the activity of carbon in the ferrite.

Fig. 10—The estimated void fraction against that measured, using
only the residual carbon concentration and the measured volume
fraction of M23C6.

4230—VOLUME 53A, DECEMBER 2022 METALLURGICAL AND MATERIALS TRANSACTIONS A



IV. CONCLUSIONS

Some really quite interesting and unexpected data
have been obtained on the tendency for hydrogen
attack, expressed as the fraction of voids generated,
have emerged from the examination of a classical
214Cr-1Mo steel subjected to a variety of heat treatments
prior to exposure.

� The heat treatment influenced the hydrogen attack
resistance. It was shown that the same steel grade
could be susceptible to or resist hydrogen attack at
the same exposure conditions, which was success-
fully correlated to the heat treatment. This suggests
that design against hydrogen attack-based solely on
temperature, hydrogen pressure and steel grade, as
done in the Nelson curves, could fall short, as was
seen in Tesoro refinery.[19]

It was possible to quantitatively measure the hydro-
gen damage by utilizing a new experimental matrix
that includes microscopy techniques combined with
synchrotron X-ray diffraction. The experimental
results helped in gaining new insights about the
mechanism such as the importance of the equilib-
rium state of carbides, which is influenced by heat
treatments. Specifying heat treatment requirements
could be a vital addition to the design standard.

� The carbon activity is not just a function of
carbide-type but also of tempering conditions that
are important in determining the tendency for
hydrogen attack. The tempering conditions influence
the equilibrium state of the carbides, hence carbon
activity and hydrogen attack susceptibility. For
example, only the heat treatment at 750 �C for 5 h
led to consistency between the thermodynamically
calculated equilibrium chromium concentration and
that measured within M23C6. Since the activity of
carbon in the ferrite depends on the Cr/Fe ratio
within the carbide, a less than expected Cr concen-
tration within M23C6 leads to a greater degree of
hydrogen attack and this is why M23C6 underwent
hydrogen attack when equilibrium state was not
reached.

� The elevated temperature tempering rate of an
initially martensitic state is slower than that of a
bainitic microstructure when considering the
sequence of substitutionally alloyed carbides. This
is probably a consequence of the greater defect
density of martensite, which means that more carbon
is retained at defects. It is well-known, for example,
that carbon located at defects has a greater stability
than in cementite, although similar studies have not
been done for alloyed carbides.

� The concentrations of carbon left in the ferrite, even
after the severe tempering treatments implemented
here, can be far in excess of the concentration in
equilibrium with M23C6. The excess has the effect of
retarding the attack of hydrogen on carbides,
because the ferrite must first be decarburised by
the hydrogen.
This somewhat contradicts a theoretical conclusion
by Schlögl et al. where it is argued that the it is

important to have as little carbon as possible in the
ferrite, based on a calculation of methane pressure
within a void.[39] But we note that only the concen-
tration of carbon in ferrite that is in equilibrium with
the carbide was considered and that those values are
much smaller than recorded here, because equilib-
rium thermodynamics does not consider the defect
structures.

ACKNOWLEDGMENTS

This work was funded by the Oil & Gas Network
Integrity Division in the Research & Development
Centre, Saudi Aramco. We acknowledge DESY (Ham-
burg, Germany), a member of the Helmholtz Associa-
tion HGF, for the provision of experimental facilities.
Parts of this research were carried out at PETRA III
and we would like to thank Dr. Magnus Hörnqvist
Colliander for assistance in using Beamline P21.2.
Beamtime was allocated for proposal I-20190135 EC.

CONFLICT OF INTEREST

The authors declare that they have no conflict of
interest.

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https://www.csb.gov/tesoro-refinery-fatal-explosion-and-fire/
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	High-Temperature Hydrogen Attack on 2.25Cr-1Mo Steel: The Roles of Residual Carbon, Initial Microstructure and Carbide Stability
	Abstract
	Introduction
	Experimental Procedure
	Material and Heat Treatment
	Hydrogen Exposure
	Characterisation

	Results and Discussion
	Carbides Precipitated Before Hydrogen Exposure
	Void Fractions After Hydrogen Exposure
	Carbide Fractions Following Hydrogen Exposure
	Effect of Deviation from Equilibrium
	Effect of Microstructure

	Conclusions
	Acknowledgments
	References