Microstructural Evolution in a 17-4 PH Stainless Steel after Aging at ...

I. INTRODUCTION
PRECIPITATION-hardened stainless steels are
widely used as structural materials for chemical and
power plants because of their balanced combination of
good mechanical properties and adequate corrosion resistance.
17-4 PH stainless steel is a martensitic stainless
steel containing approximately 3 wt pct Cu and is
strengthened by precipitation of copper in the martensite
matrix [1-8] . After a solution heat-treatment, this alloy
is precipitation hardened by tempering at about
580°C for about 4 hours. Typical service temperatures
in power plant applications are below 300°C, but increases
in hardness and tensile strength accompanied
by embrittlement was reported at temperatures ranging
from 300 to 400°C after long term aging. Since these
materials have to serve for a very long period of time
during the life span of the plants, understanding the
embrittlement mechanism at slightly above the service
temperature is very important.
The precipitation sequence in 17-4 PH stainless steel
begins with formation of coherent copper precipitates,
which occurs during the tempering treatment before service.
These coherent particles were reported to transform
to incoherent fcc-Cu particles after long term aging
at temperatures around 400°C [3] . In addition, since
the Cr concentration in 17-4 PH is within the spinodal
line, phase decomposition of the martensite into the Ferich
α and the Cr-enriched α’ is expected on aging below
450°C. Much work has shown that stainless steels
are embrittled when α’ phase precipitates by spinodal
decomposition [11] . Such α’ embrittlement is anticipated
in the 17-4 PH stainless steel as well.
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
Microstructural Evolution in a 17-4 PH Stainless Steel
after Aging at 400°C
M. MURAYAMA, Y. KATAYAMA and K. HONO
The microstructure of 17-4 PH stainless steel at various stages of heat treatment, i.e. after solution heattreatment,
tempering at 580°C and long term aging at 400°C, have been studied by atom probe field ion
microscopy (APFIM) and transmission electron microscopy (TEM). The solution treated specimen consists
largely of martensite with a small fraction of δ-ferrite. No precipitates are present in the martensite
phase, while spherical fcc-Cu particles are present in the d-ferrite. After tempering for 4 h at 580°C, coherent
Cu particles precipitate in the martensite phase. At this stage, the Cr concentration in the martensite
phase is still uniform. After 5000 h aging at 400°C, the martensite spinodaly decomposes into Fe-rich α
and Cr-enriched α’. In addition, fine particles of the G-phase (structure type D8 a , space group Fm3m )
enriched in Si, Ni and Mn have been found in intimate contact with the Cu precipitates. Following spinodal
decomposition of the martensite phase, G-phase precipitation occurs after long-term aging.
M. MURAYAMA, Researcher, and K. HONO, Head of 3rd Laboratory,
are with Materials Physics Division, National Research Institute
for Metals, Tsukuba 305-0047, Japan. Y. KATAYAMA, is
with Heavy Apparatus Engineering Laboratory, Toshiba Corporation,
Yokohama 230-0045, Japan
Manuscript submitted April 21, 1998.
Several studies on the effect of aging 17-4 PH stainless
steel were carried out [3-6] . Early work by Anthony
[4] proposed mechanical properties of 17-4 PH are influenced
by precipitation of α’ phase, but no direct evidence
for α’ precipitation was presented. Later, Jack
and Kalish [3] observed copper precipitation on aging
and correlated it to mechanical property changes; however,
there was no mention of phase decomposition in
the martensite phase. More recently, Yrieix and
Guttmann [10] reported that 17-4 PH stainless steel exhibits
high susceptibility to aging embrittlement at
400°C, and they concluded that it was essentially due
to α’ precipitation. In their study, however, no microstructural
observation results were shown. Employing
atom probe field ion microscopy (APFIM) and transmission
electron microscopy (TEM), Miller and Burke
[6] showed direct evidence for a’ precipitation after aging
at 482°C. They also reported that significant
amounts of iron, nickel and manganese were contained
in the ε-Cu precipitates even in the overaged condition.
However, their aging temperature is rather high compared
to the service condition of 17-4 PH steel.
This study aimed to carry out a more complete characterization
of microstructures in 17-4 PH stainless steel
at various stages of heat treatment, i.e., after solution
heat-treatment, tempering at 580°C for four hours, and
long term aging at 400°C, in order to obtain a better
understanding of the embrittlement phenomena on aging.
II. EXPERIMENTAL PROCEDURES
The chemical composition of the alloy used in this
study was Fe-16.5Cr-4.0Ni-3.4Cu-0.6Si-0.6Mn-0.3Nb-
0.06C (wt pct) or Fe-17.5Cr-3.8Ni-2.9Cu-1.2Si-0.6Mn-
0.2Nb-0.3C (at. pct). The alloy was solution heat-treated
at 1050°C for 1 h and subsequently water quenched.
The solution treated samples were then aged at 580°C
for 4 h (tempering). This heat treatment causes precipitation
of coherent Cu precipitates in the martensite phase

(a)
(b)
110
110
lated with the depth scale of the analysis, and the total
depth of this analysis is estimated to be approximately
60 nm. This data shows that the martensite in the solution
heat-treated specimen is a supersaturated solid solution
containing all solute atoms homogeneously.
C. Tempered microstructure
(c)
100nm
110
011
Fig. 5 (a) TEM bright field image, (b) [001] and (c) [111] SAD
pattern of the martensite phase tempered at 580°C for 4 h.
Figure 5 shows a bright field TEM image of the
martensite phase after aging for 4 h at 580°C. Brightfield
image does not give any clear contrast corresponding
to fine coherent Cu precipitates expected in this
stage. This suggests that Cu precipitates, if any, is still
coherent and they do not cause large strain contrast due
to the small strain field around the precipitate. In fact,
the SADP taken from a martensite lath does not show
any evidence for a secondary phase, suggesting that
there is no precipitates with the distinct structure different
from the bcc matrix. Figure 6 shows 3DAP elemental
mapping of Cu and Cr obtained from the martensite
phase. The Cu mapping clearly shows that there
is a small spherical particle enriched with Cu. On the
other hand, the Cr mapping shows that the distribution
of Cr atoms is uniform martensite phase. The Cu el-
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
emental map clearly shows that there is Cu enriched
precipitate in this stage. As the SADP (Figure 5(b) and
(c)) does not show any evidence for presence of the
secondary phase, the Cu enriched precipitate observe
in the 3DAP data is believed to be fully coherent bcc-
Cu. In order to quantify the concentration of the particle,
concentration depth profile was measured from
the selected region near the Cu-rich precipitate as shown
in Fig. 6 (c). The chemical composition of the Cu-rich
precipitates has been found to be 55 at.pct Cu, 30 at.pct
Fe, 10 at. pct Cr, 5 at. pct Ni. It is seen that Cr and Ni
are rejected from the Cu-enriched particle slightly. It
should also be noted that the concentration of Cu in the
particle is significantly lower than that expected from
the equilibrium e-Cu.
D. Aging for 100 h at 400°C
Figure 7 (a) and (b) show 3DAP elemental mapping
of Cu and Cr obtained from the martensite phase in the
specimen aged for 100 h at 400°C. The Cu mapping
shows that a high density of Cu-rich precipitates of approximately
3 nm in diameter is present. The Cr mapping
shows that the distribution of Cr atoms is no longer
uniform and fluctuations of Cr concentration occur.
Concentration depth profiles of Fe, Cr, Ni, Si and Cu
were measured from a selected region cutting two Curich
precipitates as shown in Figure 6(c). The concentration
of Cu in the Cu-rich precipitates is approximately
70 at. pct Cu, which is much higher than that observed
in the tempered specimen, but it is still lower than the
equilibrium concentration of the ε-Cu. The presence of
fluctuations in Cr concentration indicates that the phase
decomposition occurs in the martensite phase, and it
decomposes to Fe-rich a and Cr-enriched α’ phases. The
Cr concentration in the α’ phase is only 25 at. pct and
this is significantly lower than that of equilibrium α’
phase, suggesting that it is still in an initial stage of the
decomposition process. It should be noted that there is
no indication of partitioning of Ni.
E. Aging for 5000 h at 400°C
Figure 8 (a) shows a dark field image excited using
the 1/2(110) reflection and Figures 8 (b) – (d) shows
[001], [011], [111] zone SAD patters, respectively obtained
from the martensite phase aged for 5000 hours
at 400°C. Extra reflections indicate an ordered phase
having a cube-on-cube orientation relationship precipitates
after prolonged aging. The dark field image which
was taken using the 1/2(110) reflection shows a high
density of fine, ordered particles are dispersed in the
martensite matrix. The diffraction pattern was found to
be consistent with the G phase (structure type D8 , space a
group Fm3m ) reported in aged duplex stainless steels
[23,26,27] [19] and type 308 stainless steels .
Figure 9 shows atom probe concentration depth profiles
obtained from the martensite phase in the alloy

~26nm
(a) Cr
(b) Cu
~16nm
~26nm
aged for 5000 hours at 400°C. In addition to Cu-enriched
precipitates, significant fluctuations in Cr concentration
are observed. This shows that phase decomposition
occurs in the martensite phase, and the martensite
decomposes to Fe-rich α and Cr-enriched α’
phases. The Cr concentration in the α’ phase is approximately
40 at. pct, which is again significantly lower
than that expected from thermal equilibrium (~90 at.
pct). Thus, it is believed that this stage is still in the
middle of the decomposition process, and has not
reached the equilibrium. The Ni concentration profile
does not show any tendency of partitioning of Ni to the
a phase, unlike previous report by Danoix in the aged
ferrite phase of duplex stainless steel [26]. It appears
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
Fe & Cr / at. %
Ni
Si
Cu
100
80
60
40
20
0
30
15
0
30
15
0
60
40
20
0
0 2 4 6 8 10
Depth / ~nm
Fig. 6 3DAP elemental mapping of the martensite phase aged at 580°C for 4 h. (a) The uniform distribution of Cr and (b) a fine spherical Cu
rich precipitate is observed. (c) Concentration depth profile through a Cu precipitate.
(c)
that Cu enriched precipitates do not have any correlation
with the Cr concentration fluctuations. From the
concentration profile, two types of Cu precipitates are
observed. One is composed of Cu only, and the other is
enriched with Ni, Si and Mn, as well. The apparent Cu
concentration of the latter is ~20 at. pct Cu. This observation
is similar to the solute partitioning in an ultrafine
copper-enriched zone in a neutron irradiated A533B
submerged arc weld reported by Miller et al [13] .
Figure 10 shows 3DAP elemental mappings of Cr,
Cu, Ni and Si. Mn cannot be mapped because the mass
spectrum of Mn 2+ overlaps with that of Fe 2+ . The Cr
mapping shows the concentration of Cr fluctuates. The
Cr enriched regions appear to be interconnected, which
Fe
Cr

Fe & Cr / at. %
(a)
(b)
Ni
Si
Cu
Cr
Cu
100
80
60
40
20
0
30
15
0
30
15
0
70
50
30
10
0
~30nm
0 5 10 15 20 25
Depth / ~nm
~9nm
Fig. 7 3DAP elemental mapping of the martensite phase aged at
580°C for 4 h. (a) The phase decomposition into Cr enriched and
depleted region occur in the martensite phase. (b) Fine spherical
Cu rich precipitate. (c) Concentration depth profile obtained from
the selected region near the Cu precipitate.
is a typical feature of spinodal decomposition. The periodicity
of the fluctuation is on the order of 3 nm. The
Cu mapping shows that Cu-enriched particles approximately
8 nm in diameter are present. In direct contact
with one of the Cu particle, a Ni and Si enriched particle
is observed as indicated by an arrow in Figure
10(b). Such fine particles are always observed in contact
with Cu precipitates, and they are believed to be
the G-phase based on the TEM observation shown in
Figure 8. Figure 10 (c) shows concentration depth profiles
across the Cu precipitate and the G-phase indi-
Fe
Cr
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
Fe &Cr / at.%
(a)
(b)
110
110
(c) 011
200
(d)
100nm
Fig. 8 TEM (a) dark filed image, (b) [001], (c) [011] and (d) [111]
SAD pattern of the martensite phase aged 400°C for 5000 h. The
dark field image was taken using the 1/2(011) reflection.
Ni
Si
Mn
Cu
100
80
60
40
20
0
30
15
0
30
15
0
10
0
30
15
0
0 500 1000 1500
Number of Atoms / x50
Fig. 9 1DAP concentration depth profile of the martensite phase
aged at 400°C for 5000 h. Ni, Si, and Mn appear to be partitioned
into the Cu precipitate.
110
011

~15nm
Ni + Si Cu
(a) Cr
(b) Cu, Ni and Si
~18nm
~15nm
Fe & Cr / at. %
Ni
Si
Cu
cated by the arrow in Figure 10 (b). The Cu precipitate
contains approximately 95 at. pct Cu, which is close to
the equilibrium concentration of ε-Cu. The composition
of the G-phase is approximately 55 at. pct Ni, 25
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
100
80
60
40
20
0
100
80
60
40
20
0
50
0
100
80
60
40
20
0
Cu precipitate
0 2 4 6 8 10
Depth / ~nm
(c)
G phase
12 14
Fig. 10 3DAP elemental mapping of the martensite phase aged at 400°C for 5000 h. (a) The phase decomposition into Cr enriched and
depleted region occur in the martensite phase. (b) Large dots correspond to Cu atoms and small dots to Ni and Si atoms. (c) Concentration
depth profile across the Cu precipitate and the Ni-Si enriched particle obtained from the selected region near the Cu precipitate (indicated by
arrow).
Fe
Cr
at. pct Si, 20 at. pct Fe. In addition, Mn is also enriched
in the G-phase based on the 1DAP result shown in Figure
9 (Mn ions can be differentiated from Fe atoms using
1DAP). According to the concentration depth pro-

(a)
(b)
G phase Cu precipitate
020 fcc
011
2nm
Fig. 11 (a) HREM image taken at the [111] zone axis of the martensite
phase in the alloy aged at 400°C for 5000 h. (b) micro-diffraction
pattern obtained from the region with the Moire fringe.
file obtained by the conventional atom probe (Figure
9), Ni, Si and Mn atoms appear to be partitioned to the
Cu enriched particle indicated by the dashed lines on
the left side of the figure. However, this result is probably
an artifact caused by the convoluting effect of the
probe hole covering both Cu and G-particles because
Ni, Si and Mn enrichment is not observed at the other
Cu enriched region in Figure 9. In the latter case, the
probe hole covered only the Cu particle and missed the
G-phase which was adjacent to the Cu particle. These
results demonstrates that employment of the 3DAP technique
provides more accurate data for characterizing
the morphological feature of fine precipitates embedded
in the matrix.
Figure 11 (a) shows an HREM image taken along
the [111] zone axis of the martensite phase in the alloy
aged at 400°C for 5000 hours. In this image, fringe contrast
having a periodicity of two (011) planes is ob-
bcc
served. This is consistent with the fringe contrast expected
from the G-phase (Ni X Si , X=Fe, Mn,
16 6 7
Si, Fm3m , a=0.406 nm). Furthermore, a small particle
is observed adjacent to the G-phase, which is believed
to be a Cu precipitate. A microdiffraction pattern taken
from the region with the Moire fringes is shown in Figure
11 (b). The [111] microdiffraction pattern shows
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
the extra reflections near the {110} bcc reflections. The
δ-spacing calculated from the extra reflections is ~0.179
nm which corresponds to the spacing of the {020} fcc
planes in fcc-Cu (0.180 nm). The orientation relationship
(OR) between the particle and the martensite almost
matches with the K-S relationship.
IV. DISCUSSION
This study has clarified evolution of microstructure
in a 17-4 PH stainless steel during long term aging at
400°C. Emphasis is given to characterization of chemical
features of the precipitates which appear in the martensite
phase after prolonged aging.
The ε-Cu precipitates have been found in the δ-ferrite
after the solution heat treatment. The appearance
of these precipitates in the δ-ferrite was unaffected by
subsequent aging. Rack and Kalish [3] reported similar
microstructural feature in δ-ferrite, but they attributed
them to NbC precipitates. In this study, NbC were observed
at martensite lath boundaries with much larger
size as shown in Figure 1, and we believe that the fine
precipitate in the δ-ferrite observed in the previous study
as well as in this study are the ε-Cu. Precipitation of Cu
in the Fe-Cu binary system has been a subject of numerous
studies [14-18] , and it is well established that the
initial Cu-enriched precipitate is perfectly coherent with
the bcc matrix, while large overaged precipitates have
an fcc structure with the K-S OR. However, binary alloys
were all solution heat-treated in the single phase
ferrite region around 900°C, while the 17-4 PH stainless
steel is solution heat-treated at much higher temperature
around 1050°C. Diffusivity of Cu in the d-ferrite
at this temperature (D Cu ~ 1 x 10 -9 cm 2 /s) is more
than three orders of magnitude higher than that in the
austenite (D Cu ~ 4 x 10 -12 cm 2 /s), thus we believe that
Cu precipitated out from the δ-ferrite during cooling
after the solution heat treatment. Many Cu particles were
observed along dislocations, and this would have ease
the strain contrast for precipitation of the fcc ε-Cu. On
the other hand, no indication of presence of precipitates
are recognized in as-quenched martensite. The
solubility of Cu in the austenite phase is up to 7 at. pct
at 1100°C. However, the diffusivity of Cu in the austenite
is orders of magnitude small than that in the ferrite
phase. Thus, the precipitation kinetics are much slower
in the austenite phase and Cu can be quenched in the
martensite phase from the solution heat-treatment temperature.
After tempering at 580°C for 4 hours, fine Curich
precipitates were detected in the martensite phase
using 3DAP. The Cu content in the Cu-rich precipitates
is significantly lower than the equilibrium value for ε-
Cu. They reported that the average copper concentration
of small precipitates is approximately 50 at. pct in
the earliest stage of the precipitation. The copper concentration
in the fine, Cu-rich particles which precipitate
during tempering is in good agreement with the
early study by Goodman et al. [16] . Unlike in the d-fer-

ite, these Cu-rich precipitates observed in the martensite
phase after tempering are bcc-Cu.
The martensite phase decomposes into Fe-rich α and
Cr-enriched α′ phases by the spinodal mechanism. In
addition, fine precipitates of the G-phase has been found
in the martensite phase in direct contact with the Cu
precipitates after 5000hours aging at 400°C. Concentration
of Cr in the α′ phase is approximately 25 at. pct
after 100 hours, then reach 40 at. pct after 5000 hours
aging. This concentration is significantly lower than that
expected from the binodal line and is thought to be still
in the decomposition process. Since no precipitation of
the G-phase was observed after 100 hours aging, the
increase in yield strength up to 100 hours is attributed
to the spinodal decomposition rather than the formation
of the G-phase. G-phase precipitation was reported
in type 308 stainless steels [9, 19] and maraging steels [20]
as a grain boundary phase. In recent years, it was shown
that the G-phase exists in various Fe-Cr-Ni alloys as a
dispersed phase in the interior of the grains [21-27] . For
example, Auger et al. [27] reported phase separation and
precipitation of the G-phase in the ferrite phase of duplex
stainless steel. They concluded that the nucleation
of the G-phase takes place at the ‘interface’ between
Cr-rich a¢ and Fe-rich a. However, in the case of the
17-4 PH stainless steel, we have found that G-phase
precipitation occurs in intimate contact with ε-Cu precipitates
after the decomposition of the martensite have
progressed. This result indicates that the Cu precipitates
provides heterogeneous nucleation sites for Gphase
formation. One reason for this is that the Cu/
martensite interface is a more suitable site for nucleation
of the G-phase than the α/α’ interface. The dark
field image of the long-term aged martensite phase
shows the G-phase is dispersed uniformly in the grain
rather than grain boundary as commonly observed in
other stainless steels. This suggests that the G-phase
precipitation itself does not contribute to the
embrittlement. In fact, 80 pct of total yield strength increase
occurs before precipitation of the G-phase is
observed (100 hours aging at 400°C).
Our atom probe results have shown that the atomic
ratio of Ni to Si in the G-phase is 6:4. In addition, Mn
and Fe are also contained in the G-phase. The ternary
silicide designated as the G-phase is known to be fcc
having 116 atoms in a unit cell. The lattice parameter is
~1.12 nm. The structure of this phase is isotypic with
Th Mn (structure type D8a, space group 6 23 Fm3m ), and
the ideal composition was proposed as X Ni Si , where
6 16 7
X is typically Ti, but can be substituted for other transition
element. Considering the qualitative nature of the
atom probe result, the determined composition (55 at.
pct Ni, 25 at. pct Si, 20 at. pct Fe and some Mn) is
reasonably close to the ideal stoichiometry of X Ni Si 6 16 7
when one allows for substitution of Mn and Fe for
X [28] .
In this study, partitioning of Ni in the α phase after
Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999
spinodal decomposition of the martensite was not observed.
This result is in contrast to the APFIM results
by Danoix et al. [26] which reported that nickel is rejected
from Cr-enriched α’ phase and partitioned into
the Fe-rich α phase in the aged ferrite phase of duplex
stainless steel. In the equilibrium condition, the solubility
limit of Ni in α and α’ were estimated to be 4.1
and 0.01 at. pct, respectively, by the ThermoCalc software.
Thus, Ni should partition into the α’phase in 17-
4 PH as well. This suggests Ni partitioning does not
occur until decomposition progresses further and the
driving force for the partitioning reaction increases.
Similar delayed partitioning behavior was reported for
Al in an Fe-Cr based alloy [29] . As the Ni concentration
of the duplex stainless steel was higher than that in 17-
4 PH, the driving force for Ni partitioning is much higher
in the duplex stainless steel. Thus the absence of concurrent
partitioning of Ni with spinodal decomposition
in 17-4 PH is reasonable.
V. CONCLUSIONS
The microstructural features of 17-4 PH stainless
steel at various stages of heat-treatment has been investigated
by APFIM and TEM. The main results are
as follows:
1. The fcc Cu-rich particles precipitate in the δ-ferrite
grains during cooling after the solution heat treatment.
The precipitates have the K-S orientation relationship
with the δ-ferrite matrix.
2. Tempering at 580°C for 4 hours results in precipitation
of coherent bcc particles in the martensite phase.
The composition of the precipitates is approximately
60 at. pct Cu and Cr, Ni and Si are rejected from the
bcc-Cu. The Cr concentration is homogeneous in
the martensite phase at this stage.
3. After 100 hours aging at 400°C, evidence for
spinodal decomposition of the martensite phase into
Fe-rich a and Cr-enriched a’ phase is found. The Cr
concentration in the a’ phase is approximately 25 at.
pct, significantly lower than the equilibrium value.
The concentration of Cu in the Cu-enriched precipitate
is approximately 70 at. pct at this stage.
4. Spinodal decomposition of the martensite phase
progresses further after 5000 hours aging. The Cr
concentration of the α’ phase is 40 at. pct at this
stage. The structure of the coarsened Cu precipitates
is fcc and their concentration is almost 100 at. pct
Cu. Fine particles of G-phase containing approximately
60 at. pct Ni, 25 at. pct Si and some Fe and
Mn are present in direct contact with the Cu precipitates,
suggesting that G-phase is heterogeneously
nucleated at the martensite/Cu interface.
5. The increase in hardness and yield strength of the
17-4 PH stainless steel after aging at 400°C is mostly
caused by spinodal decomposition of the martensite
phase. G-phase precipitation does not appear to make

a significant contribution to the embrittlement during
aging.
ACKNOWLEDGMENTS
The authors thank Professor W.T. Reynolds, Jr., Virginia
Polytechnic Institute and State University, for
valuable discussions. This work was supported by the
Frontier Research Center for Structural Materials,
NRIM.
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