Available online at www.sciencedirect.com
Corrosion Science 50 (2008) 1481–1491
www.elsevier.com/locate/corsci
The use of SIMS in quality control and failure analysis
of electrodeposited items inspected for hydrogen effects
E. Kossoy a, Y. Khoptiar a, C. Cytermann b, G. Shemesh a, H. Katz a,
H. Sheinkopf a, I. Cohen a, N. Eliaz c,*
a
Materials Division, Depot 22, Israel Air Force, P.O. Box 02538, Israel
Solid-State Institute, Technion – Israel Institute of Technology, Haifa 32000, Israel
c
Biomaterials and Corrosion Laboratory, School of Mechanical Engineering, Tel-Aviv University, Ramat Aviv 69978, Israel
b
Received 18 December 2007; accepted 9 January 2008
Available online 16 February 2008
Abstract
In high-strength steels it is often difficult to distinguish between hydrogen embrittlement and various other brittle failure mechanisms.
The objective of this work was to develop a sensitive analytical procedure based on secondary ion mass spectrometry (SIMS) that would
allow in-service identification of local hydrogen accumulation, either during quality control or during failure analysis of electroplated
items. Dynamic SIMS was found useful in identifying when baking of Cd-plated AISI 4340 steel was not carried out, thus potentially
leading to hydrogen embrittlement. In all non-baked samples, an increase in the hydrogen signal was found at the Cd/steel interface. In
baked samples, either a peak was not observed at the interface, or it was insignificant based on determination of the ratios between the
hydrogen signals in the coating, interface and substrate. This reproducible effect was monitored even after 16 months storage in a desiccator. These observations make the procedure practical in suggesting more accurate, reliable and cost effective recommendations for
prevention of failures. The main effect of baking was found to be effusion of hydrogen from the interface and the substrate steel into the
atmosphere. A mechanism for delayed failure is suggested.
Ó 2008 Elsevier Ltd. All rights reserved.
Keywords: A. Steel; B. SIMS; C. Electrodeposited films; C. Hydrogen embrittlement; Failure analysis
1. Introduction
Failure analysis is usually a necessary stage in determining the mechanism/s and cause/s of failure, so that effective
corrective actions can be implemented and recurrences of
the failure can be eliminated (or, at least, minimized) [1].
It is well known, however, that in the case of high-strength
steels it is often difficult to distinguish between various brittle failure mechanisms because similar fracture characteristics may appear in each of them [2]. In particular, it is
sometimes nearly impossible to tell apart hydrogen embrittlement (HE) and stress corrosion cracking (SCC) in prac-
*
Corresponding author. Tel.: +972 3 640 7384; fax: +972 3 640 7617.
E-mail address: neliaz@eng.tau.ac.il (N. Eliaz).
0010-938X/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved.
doi:10.1016/j.corsci.2008.01.016
tical failure situations [3]. Obviously, the recommendations
for prevention of failure recurrence would be totally different in each case.
High-strength steels might suffer severe loss of ductility,
toughness and strength due to HE [4], resulting in sudden
fracture (i.e., delayed failure). The susceptibility of steels
to HE increases with tensile strength. The embrittlement
can be caused by either external or internal hydrogen.
Hydrogen can be introduced into the metal either during
fabrication (e.g. during cleaning, pickling, phosphating,
electroplating, autocatalytic processes, welding or brazing
operations, as a result of lubricant breakdown, etc.) or in
service (mainly, due to cathodic protection reactions or
corrosion reactions) [5]. Atomic hydrogen is codeposited
and absorbed in the metal substrate during electroplating
[6,7]. In coating systems with low Faradaic efficiency
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E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
(FE), hydrogen evolution during electrodeposition is
enhanced, thus increasing the likelihood of HE. For example, hard chromium coated items are more prone to HE
than those coated with cadmium [2]. Even the high efficiency etch-free (HEEF) hard chromium coating, in which
the FE was increased from 15% that is typical of conventional hard chromium to 25%, has been found susceptible
to HE [8]. In highly efficient systems, such as Cu and Ag
baths, codeposition of hydrogen occurs only when the limiting current density is exceeded, or when the added complexing agents shift the potential of metal deposition to
sufficiently negative values. The higher the hydrogen overpotential on a given metal, the lower the amount of hydrogen absorbed in it [6]. One of the mechanisms that have
been suggested to explain the HE of steels is the high-pressure bubble formation [4]. Eliaz et al. [9] have recently
developed a coupled diffusion/fracture mechanics
approach to describe the expansion of high-pressure hydrogen bubbles and propagation of cracks between them in the
absence of external loads, allowing determination of the
time to failure.
Heat treatment (‘‘baking”) is commonly employed following electroplating of various coatings in order to render
the normally mobile hydrogen immobile. For high-strength
steels electrodeposited with hard chromium, for example,
this is typically done at 177–205°C for at least 3 h. The
treatment is required for all steel parts hardened to 40 RC
or more, and should be applied not later than 4 h after
the completion of the plating process [10]. Evaluation of
the effect of residual hydrogen on the mechanical properties is then done according to standards such as ASTM F
519 [11] and ASTM F 1624 [12]. Certain very susceptible
alloys require the use of fabrication processes other than
electroplating.
Although hydrogen diffusion is exponentially increased
with temperature, the required time to reduce the diffusible
hydrogen concentration to a given level increases with the
square of the section thickness. For thick sections this
can mean hundreds of hours at any common baking temperature. Even then, there is no guarantee that permanent
damage or irreversible HE has not already occurred [2]. In
the case of Cd plating, for example, the high solubility of
hydrogen in Cd relative to that in the iron-based substrate
makes the Cd-layer a source during the initial stage of baking; consequently, the hydrogen concentration in the steel
substrate may increase [13]. Moreover, because the diffusivity of hydrogen in Cd is much lower than that in steel, the
Cd-layer acts as a diffusion barrier to outgasing of hydrogen during baking, and a significant concentration of dissolved hydrogen might remain in the steel, even after
baking times of 100 h [14].
If standard procedures are followed precisely, HErelated failures of electroplated items can usually be prevented. However, if improper baking process was carried
out, or such a treatment was not carried at all, delayed failure might occur. In many practical cases of in-service failures, even when a failed item is suspected for HE, it is not
easy to confirm whether all manufacturing processes were
carried out properly. Then, in order to prevent failures of
items from the same manufacturing batch, there is often
no choice but to remove the entire batch from service. This
solution is usually complicated and costly, in particular
when considering critical aeronautical components. Furthermore, taking into account that random hydrogen
charging often cannot be eliminated or easily controlled
by platers, testing representative quantities of the finished
items is necessary in the framework of quality control [5].
Hence, an accurate, sensitive experimental procedure that
is capable of mapping the level of residual hydrogen would
be very useful both at the stage of quality control and at the
stage of failure analysis. This was the goal of the present
work. Because high-strength steels might be embrittled
even by very low concentrations of hydrogen, and because
the ability to reach locally a critical concentration of
hydrogen is more significant than the mean concentration
of hydrogen in the bulk of the metal, local hydrogen quantification is necessary for practical evaluation of the risk of
HE. Experimental methods such as the hydrogen microprint technique [15,16], tritium autoradiography [17,18]
and thermal desorption spectroscopy (TDS) [19] were evaluated and found unsuitable to accomplish this project’s
goal.
In secondary ion mass spectrometry (SIMS), a solid
sample material is bombarded with a focused primary ion
beam (300 eV–30 keV). The primary ions are implanted
into the sample down to tens of nm. A transfer of kinetic
energy takes place between the surface atoms and collision
cascades are created. Some collisions are oriented backwards toward the surface. If they have enough energy to
overcome the surface barrier potential, atoms, molecules
and molecular fragments are ejected (sputtered) from the
sample surface. Most of these particles are neutrals, but a
small fraction is ionized, either positively or negatively.
These secondary ions, characteristic of the composition
of the analyzed volume, are separated according to their
mass-to-charge ratio and collected. SIMS combines simultaneous detection of ions over a virtually unlimited mass
range (from hydrogen to uranium), at high mass resolution
(up to 0.02 amu), at sensitivities that may be as high as
ppb-ppm, with mapping of the lateral distribution of species and with depth profiling. Its drawbacks include the
requirement for UHV-compatible and size-limited samples,
a very high sensitivity for surface morphology and surface
contaminants, typically limited optical capabilities that
make it difficult to find grains or local regions of interest
for analysis, the need for an expert operator, and a relatively high cost [20].
There are two major types of SIMS instruments: static
and dynamic. In static SIMS, a very low ion beam current
density is used and less than 1% of the original surface of
the sample is consumed during the course of the analysis.
Its time-of-flight (TOF) mass analyzer separates the ions
at a field-free drift path according to their kinetic energy.
The analysis and the sputtering (if conducted) are con-
E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
ducted separately. The spatial resolution is 0.1 lm (imaging), the depth resolution is several angstroms and, because
of the small primary intensities used, the sputtering depth is
usually limited to 1lm. Static SIMS can suggest molecular structures and observe extremely high mass fragments
(e.g. of polymers), which dynamic SIMS cannot do. In
dynamic SIMS, a much higher primary ion beam current
density is used. The surface of the sample is continuously
etched away during the course of the analysis. A quadrupole or magnetic sector mass analyzer separates the masses,
where only masses of choice are able to pass. The spatial
resolution is a few micrometers (imaging), the depth resolution is a few nanometers (depending on the primary ion
energy), and a sputtering depth of tens of micrometers is
possible [20].
SIMS has already been applied to a limited extent in
mechanistic investigations of HE and SCC, with some success, as well as to identify the traps for hydrogen isotopes
in different alloys. Takai et al. [21,22] used SIMS to visualize the trapping sites of deuterium in high-strength steels.
Line scans showed that the concentration of deuterium at
inclusions, grain boundaries, and phosphorous segregation
bands in a steel bar for prestressed concrete was 11.0, 7.8
and 5.0 times higher than that in the matrix, respectively.
Fractured samples were analyzed following FIP test for
hydrogen-induced SCC. Based on TDS analysis it was concluded that it is necessary to promptly carry out SIMS
analysis of the trapped deuterium, either around inclusions
or at the grain boundaries, within 24 h after the delayed
fracture test, otherwise the measured signal would be too
low [21]. Later [23], these authors combined SIMS with
TDS in order to visualize the initial deuterium distributions
in a spheroidal graphite cast iron, and to relate the desorption profiles during different heat treatments with trapping
sites. Deuterium was used instead of hydrogen as the
detected ion in order to increase the sensitivity. Brass
et al. [24] performed tritium autoradiography and deuterium profiling by SIMS on pre-cracked and then boltloaded double cantilever beam (DCB) specimens made of
4120 HSLA steel in order to measure the hydrogen concentration in the plastic zone at the crack tip. The samples
were cathodically charged, repolished and quenched at
196°C in order to prevent deuterium desorption. A three
orders of magnitude drop in the deuterium concentration
was observed after 100 h aging at 20°C. Next, Chene
et al. [25] introduced deuterium into flat tensile specimens
made of Ni-based alloy 600 by potentiostatic cathodic
polarization. SIMS profiles were calibrated by thermal
extraction of deuterium and correlated to fracture surfaces
in order to assess the effect of deuterium concentration on
the extent of intergranular fracture pattern. Sastri and
McDonnell [26] applied the hot-extraction thermal conductivity method, SIMS, anodic dissolution, and the reaction
of the liberated hydrogen with nitroxyl free radical method,
in the determination of surface enriched hydrogen in AISI
1062 and 4037 steels. The results supported the hypothesis
that HE originates in the thin outer surface layers of these
1483
steels. Mao and Li [27] used SIMS to measure the hydrogen
distribution around a crack tip in compact tensile (CT)
specimens made of pipeline steel X-80. These specimens
had been loaded by a wedge and subjected for 72 h to
SCC. The results indicated generation and accumulation
of hydrogen around the SCC crack tip. Shvachko [28] used
the SIMS to investigate the hydrogen state on pure iron
pre-charged electrolytically with hydrogen. It was found
that the ability to measure negative ion emission extends
the analytical capability of SIMS in the investigation of
Fe–H interaction phenomena, and that the hydrogen
atoms diffusing from the bulk of iron can bear a negative
charge on its surface.
None of references [21–28] deals with, or is directly
related to, electroplating, baking, quality control and failure analysis of real, engineering parts. They all involved
analysis of deuterium (instead of hydrogen), polished samples, quenching at cryogenic temperatures to prevent
hydrogen desorption before SIMS analysis, or performance
of the SIMS analysis quickly after the hydrogen was introduced into the material and/or delayed failure occurred.
These characteristics, in addition to the use of specimens
that are not always easy to extract from in-service parts,
make them less useful for the type of tasks of interest in this
study. Therefore, the objective of the present research was
to develop an in-service analytical procedure based on
SIMS that would allow identification of local hydrogen
accumulation either during quality control or during failure analysis of electroplated items. This procedure should
be helpful in distinguishing between HE and other possible
failure mechanisms in high-strength steels. Availability of
such a method would give way to more accurate, reliable
and cost effective recommendations for the prevention of
failure recurrence. It could also be used in courts, when
the aptness of the baking process is questioned. To the best
of the authors’ knowledge, such use of SIMS has not been
reported before elsewhere.
For this work, AISI 4340 low alloy steel was selected.
This steel combines deep hardenability with high ductility,
toughness and strength. It is typically used in heavy-duty
structural parts, such as aircraft landing gears, power
transmission gears and shafts, etc. This steel becomes susceptible to HE when heat treated to tensile strengths above
1400 MPa (200 kpsi). Hence, parts exposed to hydrogen
during manufacturing are usually baked subsequently. In
addition, this steel exhibits extremely poor resistance to
SCC when tempered to tensile strengths of 1500–
1950 MPa (220–280 kpsi).
2. Experimental procedures
2.1. Dynamic SIMS analyses
At an early stage of the project, samples were analyzed
by static SIMS. However, due to several limitations that
were identified, it was eventually decided to focus on
dynamic SIMS.
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E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
Fig. 1. A standard [10] tensile V-notch (TVN) sample made of AISI 4340 steel heat treated to 50–53 RC.
2.1.1. Instrument specifications, calibration and operating
conditions
A Cameca ims 4f dynamic SIMS instrument was used. A
18–25 nA, 5.5 keV Cs+ primary ion beam was rastered over
a sputtering square area that varied between 50 50 lm
and 250 250 lm, while the analyzed area was set to
33 lm (in diameter) at the center of the crater. The secondary negative ions H , Fe , C , O and CdO were monitored. The CdO– secondary ions was preferred over Cd–
because the latter has a very low secondary yield. The samples were biased at 4.5 kV. The choice of H– is related to
its higher sensitivity in SIMS. The samples were introduced
into the SIMS chamber at least 12 h before the measurement in order to minimize the environment contribution
to the hydrogen secondary yield. The base pressure before
analysis was 2 10 10 Torr. The resulting crater depths
were measured with the aid of a depth profilometer (model
Tencor 200), with a precision of 5%.
An attempt was made to calibrate the instrument for
quantitative analyses of the hydrogen concentration. In
addition, the sputtering rate was determined for the AISI
4340 steel. To these aims, a sample with dimensions of
7 12 5 mm was cut from a steel bar hardened to 50–
53 RC. The sample was polished in a water-based alumina
suspension to 0.05 lm, cleaned ultrasonically in 2-propanol
for 1 min, air dried, and stored in a desiccator. Next,
hydrogen ions implantation was carried out into 2/3 of
the area, at a dose of 1015 atoms/cm2 and ions energy of
40 keV. Computer simulation predicted that under these
conditions, a maximum hydrogen concentration of
7 1019 atoms/cm3 would be established at a depth of
216 ± 61 nm. After ions implantation, the specimen was
kept in liquid nitrogen until the SIMS analysis began. That
analysis, however, did not reveal any difference in the
hydrogen concentration within the implanted and nonimplanted areas. This could be either due to a high background contribution from hydrogen in the vacuum chamber or, more likely, due to rapid desorption of hydrogen
into the vacuum phase (the equation of diffusion length
may be used to estimate the time required for the escape
of hydrogen from the thin ion-implanted layer). By measuring the crater depth after sputtering, the typical sputtering rate was determined. Thus, it was found that the
sputtering rate was increased from 0.18 nm/s to 1.2 nm/s
and then to 3.2 nm/s as the sputtering area was reduced
from 250 250 lm to
50 50 lm, respectively.
100 100 lm
and
then
to
2.1.2. Analysis of a fracture surface
Since local accumulation of hydrogen generally leads to
failure initiation by intergranular cracking (IGC), a correlation was sought between the local hydrogen signal and
the fracture pattern at that analysis point. A standard
[11] tensile V-notch (TVN) specimen (Fig. 1), heat treated
to hardness of 50–53 RC, was electroplated with Cr in
accordance with QQ-C-320 [10]. No baking was performed
afterwards. The reason for selecting Cr and not Cd was to
maximize the effect of hydrogen on the mechanical properties and fracture mode. When the specimen was loaded
statically to 77% of the steel ultimate tensile strength, it
fractured within a few seconds. Two cylindrical samples,
each approximately 1 cm long and containing one of the
opposing fracture surfaces, were cut. The fracture surface
was first examined under a stereoscope. Next, one fracture
surface was analyzed using a Scanning Electron Microscope (Jeol JSM-7000F Field-Emission SEM). The other
sample was cleaned ultrasonically for 1 min in 2-propanol,
air dried, and kept in a desiccator for 46 days before SIMS
analysis was carried out. This long storage period after
fracture was chosen as to mimic real failure analysis events.
The depths of measurement were between 3.1 and 4.5 lm.
2.1.3. Analyses through Cd plating
A second approach was to measure hydrogen depth profiles during sputtering through a Cd electroplate into the
substrate steel. Since the electroplating protects the metal
surface from contamination and prevents outgassing of
hydrogen, this approach was expected to provide more reliable results. It may be useful both in quality control and in
failure analysis, for instance by acquiring depth profiles
close and in parallel to the fracture surface. A comparison
was made between baked and non-baked samples.
Two rectangular samples were prepared from steel hardened to 40–42 RC. The samples were ground on a 120-grit
paper and cleaned ultrasonically for 1 min in 2-propanol.
Following cleaning in an alkaline bath, cadmium electroplating was applied for 1.5 min, in accordance with QQP-416, type I, class 3 [29]. One sample, designated herein
as B1, was immediately baked at 191°C for 23 h. The other
sample, designated herein as NB1, was not baked. The
E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
thickness of the coating was measured on metallographic
cross-sections by means of a Reichart–Jung MeF-3 light
microscope. It was found to be 5 lm. The remaining parts
of the two samples, about 6 6.5 4 mm each, were
cleaned ultrasonically for 1 min in 2-propanol, air dried,
and kept in a desiccator until SIMS analysis started. Overall, nine days passed from electroplating to measurement.
The sputtering area was 250 250 lm, whereas the depths
of measurement were between 5.0 and 9.5 lm.
In order to demonstrate reproducibility, seven 9 6
4.5 mm samples were prepared as described above, except
for increasing the plating time to 4 min. Three samples,
designated herein as B2, B3 and B4, were immediately
baked at 191°C for 24 h. Three samples, designated herein
as NB2, NB3 and NB4 were not baked. The seventh sample was not baked, but served for measuring the coating
thickness. The latter was found to be 10–15 lm. The six
samples for SIMS analysis were stored in a desiccator for
56 d before SIMS analysis. SIMS depth profiles were
acquired at three different points on each sample. In each
measurement, which lasted about 20 min, the primary ion
1485
beam current was 20 nA, the analyzed area was 33 lm in
diameter, and the sputtering area was 125 125 lm. The
measurements were terminated when all signals became
stable, within the bulk of the steel.
Finally, in order to confirm whether any observed phenomenon can be identified even after extremely long periods, simulating the worse scenario in real-life failure
analyses, two Cd-coated samples were arbitrarily chosen
and reanalyzed after 16 months storage in a desiccator.
One sample (designated herein as B5) was originally baked
immediately after electroplating, while the other one (designated herein as NB5) was not.
3. Results and discussion
3.1. Analysis of a fracture surface
The high-strength TVN specimen that was electroplated
with Cr without performing subsequent baking (see Section 2.1.2) exhibited a brittle failure mode. Fractography
(Fig. 2) revealed two distinct regions. A peripheral
Fig. 2. SEM images of the fracture surface of a Cr-plated, non-baked TVN sample, fractured at 77% of the ultimate tensile strength. The fracture surface
consists of a peripheral intergranular region (a) and an inner dimpled (overload) region (b).
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E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
intergranular fracture surface was found to exist along
roughly half of the perimeter, penetrating approximately
120 lm into the sample. Such an intergranular pattern
along prior austenite grain boundaries is typical in failures
of high-strength steels due to HE. The rest of the fracture
was characterized by a dimpled rupture, typical of a failure
by overload, with randomly oriented 50–100 lm long
cracks.
SIMS analyses were made at four points, two within the
IG region and two at the center of the dimpled area. Note
that the size of grains, which varied between 10 and 30 lm,
was comparable with the analyzed area (33 lm in diameter). No correlation between the intensity of the hydrogen
signal and the fracture pattern was observed. The reasons
for this inability to map a hydrogen-enriched region may
be: (1) A rapid escape of hydrogen from the fracture surface due to its high diffusivity in the steel, (2) the high
roughness of (any) fracture surface, (3) adsorption of contaminants on the fracture surface, masking the hydrogen
effect, and (4) a too high hydrogen background signal.
Thus, it was decided to evaluate an alternative approach
for identifying the residual hydrogen.
3.2. Analysis through Cd plating
In contrast to the results discussed so far, hydrogen
depth profiles through the Cd plating revealed clear differences between the baked and the non-baked samples. Typical logarithmic plots of the secondary ion yields as a
function of time are presented in Figs. 3a and b for the
baked and non-baked samples, respectively. It should be
noted that the sputter rate in the Cd plating is much higher
than in the underneath steel. Therefore, to plot the secondary ion yield as a function of depth would have required a
careful point weighted calculation of the sputter rate at the
Cd/steel interface. In addition, in order to compare
between the measurements, the profiles were normalized
by the average Fe– intensity in the steel. A sharp Cd/steel
interface can easily be distinguished in Fig. 3, despite the
roughly ground (namely, 120 grit) surface. The interface
is characterized by a decrease in the signal of CdO and
an increase in the signals of Fe and C , when penetrating
into the steel substrate. Thus, it may be concluded that
SIMS depth profiles of CdO , Fe and C may be used
to identify the Cd/steel interface. The initial decrease in
the hydrogen ion signal (as well as in other ion signals)
may be attributed to the implantation and accumulation
of Cs+ ions from the sputtering ion beam, until saturation
is attained. It has already been reported that, while being
accumulated in the sample, Cs+ ions may affect the extent
to which surface atoms are ionized [21].
The most important observation in this study was that
the hydrogen signal was increased at the Cd/steel interface
in all non-baked samples. This phenomenon is illustrated
in Fig. 3c, in which the intensity is drawn on a linear
scale. Comparison of the hydrogen profiles for the baked
and for the non-baked samples enabled to exclude the
attribution of variation in the hydrogen level to changes
in the hydrogen ionization yield in different media
(namely, steel versus Cd). Thus, a simple depth profile
through the plating may aid in determining when HE
due to internal hydrogen might become a concern in an
electroplated item. SIMS may be a very sensitive and reliable technique for identifying when baking was not done
at all, or was done improperly. From the fracture pattern
shown in Fig. 2 it could be estimated that hydrogen accumulation and damage was most significant around the
coating/steel interface, where intergranular fracture pattern was observed. Results from tensile tests, either sustained load tests or incremental step load tests, have
been used before also by other researchers to indirectly
support the mechanism according to which baking after
plating allows the atomic hydrogen to be removed and
effuse through the coating and also be more homogeneously distributed in the bulk of the substrate metal.
Paatsch and Hodoroaba [30] have applied the glow discharge emission optical spectroscopy (GD-OES) technique and showed that hydrogen depth profiles may be
used to distinguish between baked and non-baked samples
in systems such as Zn on steel, Cd on steel and Co–Pt–W
on Cu [30]. The exact concentration of hydrogen could
not be determined in that study.
Four questions may be raised at this stage: (1) Is the
aforementioned effect observed by SIMS in non-baked
samples reproducible? (2) May it appear also in properly
baked samples? (3) How sensitive is it to the location of
SIMS measurement along the sample? (4) Is it possible to
observe it even after longer storage periods? In order to
answer questions (1) through (3), six more samples were
analyzed – three with baking and three without (see Section
2.1.3). SIMS depth profiles were recorded at three different
positions in each sample. In all measurements, a peak in
the signal of hydrogen was observed at the Cd/steel interface in non-baked samples. This is demonstrated in
Fig. 4c for sample NB3. Therefore, it may be concluded
that the detected effect seems to be reproducible. Of course,
it is advisable to carry out more experiments in order to
increase the population and provide a proper statistical
analysis.
A short explanation of the hydrogen background should
be provided at this stage. The hydrogen background
increases as either the raster size is increased or the sputtering rate is decreased. Hence, the background for all samples presented in Fig. 4 must be relatively low. The
background is different in Cd and in steel because of the
difference in sputtering rate (in the Cd-layer it is unlikely
to be high because of the high sputtering rate for Cd). This
was confirmed in baked samples, in which a decrease of the
raster size from 150 to 50 lm did not result in further
decrease in the hydrogen signal in the steel substrate. The
implication of this observation is that the hydrogen signal in the steel is meaningful at a raster size of 125 lm,
which was used when acquiring the profiles presented in
Fig. 4.
E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
1487
Fig. 3. Typical dynamic SIMS depth profiles measured across the Cd plating in: (a) a baked sample, (b) a non-baked sample, and (c) comparison of typical
hydrogen depth profiles in baked and in non-baked samples, with the intensity being drawn on a linear scale.
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E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
Fig. 4. Hydrogen (H ) depth profiles obtained from two baked (a–b) and one non-baked (c) samples. In each sample, the profile was acquired at three
different regions.
E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
Table 1
Quantitative characteristics of the H depth profiles in baked (B) and nonbaked (NB) Cd-plated samples
Sample name–
number of
measurement
Hinterface =Hcoating
Hsteel =Hcoating
Hinterface =Hsteel
B2–1
B2–2
B2–3
B3–1
B3–2
B3–3
B4–1
B4–2
B4–3
1.00
1.28
1.20
1.00
1.00
1.00
1.38
1.00
1.50
0.125
0.170
0.175
0.107
0.085
0.132
0.519
0.066
0.300
8.00
7.53
6.86
9.35
11.76
7.58
2.66
15.15
5.00
NB2–1
NB2–2
NB2–3
NB3–1
NB3–2
NB3–3
NB4–1
NB4–2
NB4-3
2.59
3.11
2.53
6.54
2.50
9.45
4.70
2.51
4.30
1.310
0.791
0.683
1.610
0.570
4.450
2.090
0.807
1.500
1.98
3.93
3.70
4.06
4.39
2.12
2.25
3.11
2.87
Hinterface is the strongest signal detected at the Cd/steel interface; Hcoating is
the weakest signal within the coating layer, closest to the interface; Hsteel is
the signal at the deepest point analyzed within the substrate steel.
Referring to question (2), nine hydrogen depth profiles
were recorded from samples B2–B4. In five of these, no
hydrogen peak was observed at the Cd/Steel interface
[see, for example, Fig. 4a for sample B3]. However, in the
remaining four profiles, a small hydrogen peak was
observed at the interface [see Fig. 4b for sample B2]. Fortunately, a complementary quantitative analysis can be
suggested (see below) in order to prevent an unambiguous interpretation of SIMS data. To the best of our knowledge, this is the first time ever that such an analysis is
proposed.
Table 1 summarizes three quantitative characteristics of
the H depth profiles in baked and non-baked Cd-plated
samples, namely the Hinterface =Hcoating ; Hsteel =Hcoating
and Hinterface =Hsteel ratios. Here, Hinterface is the strongest signal detected at the Cd/steel interface, Hcoating is the weakest
signal within the coating layer – closest to the interface, and
Hsteel is the signal at the deepest point analyzed within the
substrate steel. In all cases, the depths of measurement
were between 6.14 and 9.63 lm. The values below are presented in terms of (mean ± standard error). From Table 1
it is evident that while the Hinterface =Hcoating ratio was 1.00–
1.50 (1.15 ± 0.07) in the baked samples, it increased to
2.50–9.45 (4.25 ± 0.80) in the non-baked samples. In addition, while the Hsteel =Hcoating ratio was 0.066–0.519 (0.19 ±
0.05) in the baked samples, it increased to 0.570–4.450
(1.53 ± 0.40) in the non-baked samples. Finally, while the
Hinterface =Hsteel was 2.66–15.15 (8.21 ± 1.21) in the baked
samples, it reduced to 1.98–4.39 (3.16 ± 0.30) in the nonbaked samples. In order to determine whether the differences between the means for the baked and for the non-
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baked samples are statistically significant, a Student’s t-test
was performed, using Microsoft Excel software. For a
probability level p = 0.05 (i.e. 95% probability of making
a correct statement) and 16 degrees of freedom
(n1 + n2 2), the tabulated t-value equals 2.12. The calculated t-values were 3.88, 3.33 and 4.04 for Hinterface =Hcoating ;
Hsteel =Hcoating and Hinterface =Hsteel , respectively. Because all
three calculated t-values are higher than the tabulated
value, it may be concluded that the means are significantly
different for the baked and for the non-baked samples.
Thus, it may be concluded that baking resulted in a significant decrease in the concentration of hydrogen in the steel
compared to its concentration in the Cd coating. In addition, while following baking the Cd/steel interface became
less enriched in hydrogen compared to the Cd coating
itself, the steel substrate became less enriched in hydrogen
compared to the interface. This suggests that the main
effect of baking was effusion of hydrogen from the substrate steel and the Cd/steel interface, through the coating,
into the atmosphere. Yet, it cannot be excluded that
homogenization of hydrogen distribution within the bulk
of the substrate steel also occurred to some extent as a
result of hydrogen diffusion from the interface zone deeper
into the steel, where SIMS analysis was not carried out.
It is important to mention that both the iron and the
oxygen depth profiles did not differ between the baked
and the non-baked samples. Thus, it can be concluded that
the changes in the intensity of the hydrogen signal reflected
changes in the amount of hydrogen, rather than changes in
its ionization yield due to changes in the yield of other ions.
Furthermore, referring to question (3) above, it seems that
the same effect may be identified at different locations along
the surface of the samples; so, it does not limit the investigator to a specific zone, which might be more difficult to
track in real in-service failure analyses.
Finally, referring to question (4), two samples – one
baked and one non-based – were reanalyzed by dynamic
SIMS 16 months after plating. Remarkably, in contrast
to all previous publications (see, for example, Refs.
[21,24]), a similar effect was monitored, as evident in
Fig. 5. This finding has two important implications. Firstly,
it indicates that HE-related delayed failures of improperly
baked electroplated items may be related to the time-independent accumulation of hydrogen at the coating/substrate
interface, and not necessarily to irreversible damage that
occurred in the substrate metal during fabrication – as
usually claimed in the literature. Depending on the concentration of hydrogen at the interface, blistering and delamination might occur, or the interface may serve as a
reversible hydrogen trap that provides a reservoir of diffusible (mobile) hydrogen for the steel. In the latter case, in an
event that energy is absorbed in the material which is sufficient to overcome the activation energy barrier for hydrogen detrapping from the interface, so that diffusion of
hydrogen towards strong irreversible traps in the steel is
enhanced, the base metal itself may become embrittled
within its bulk. Secondly, it seems that there is no time con-
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E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
Fig. 5. Typical dynamic SIMS depth profiles measured after 16 months storage in a desiccator across the Cd plating in: (a) a baked sample and (b) a nonbaked sample.
straint for the use of SIMS analyses in failure analyses of
electroplated items suspected for hydrogen damage.
4. Conclusions
This paper demonstrated the effectiveness of dynamic
SIMS in recognizing improper baking of Cd-electroplated
AISI 4340 steel, which resulted in hydrogen embrittlement.
The following conclusions may be drawn:
(1) In non-baked samples, an increase in the hydrogen
signal is found at the Cd/steel interface. In baked
samples, either a peak is not observed at the interface,
or it could be ignored based on determination of the
ratios between the hydrogen signals in the coating,
interface and substrate.
(2) The main effect of baking seems to be effusion of
hydrogen from the interface and the substrate steel
into the atmosphere.
(3) The reproducible effect can be monitored even after
very long storage times (e.g. 16 months storage in a
desiccator).
(4) HE-related delayed failures may be explained in
terms of the time-independent reservoir of hydrogen
at the coating/substrate interface, rather than in
terms of irreversible damage that occurred within
the substrate during electroplating.
(5) The suggested procedure may be used either during
quality control or during failure analysis of electroplated items. It should be practical in suggesting more
accurate, reliable and cost effective recommendations
for prevention of failures (e.g. by distinguishing
between HE- and SCC-related failures).
E. Kossoy et al. / Corrosion Science 50 (2008) 1481–1491
Acknowledgements
The authors are grateful for the financial support of the
Israel Ministry of Defense. The authors thank A. Gladkich
for conducting the static SIMS tests.
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