Download - Microstructural Evolution in a 17-4 PH Stainless Steel

Published in Metall. Mater. Trans. A. Vol. 30A, pp. 345-353. 1999

I. INTRODUCTION

PRECIPITATION-hardened stainless steels arewidely used as structural materials for chemical andpower plants because of their balanced combination ofgood mechanical properties and adequate corrosion re-sistance. 17-4 PH stainless steel is a martensitic stain-less steel containing approximately 3 wt pct Cu and isstrengthened by precipitation of copper in the marten-site matrix [1-8]. After a solution heat-treatment, this al-loy is precipitation hardened by tempering at about580°C for about 4 hours. Typical service temperaturesin power plant applications are below 300°C, but in-creases in hardness and tensile strength accompaniedby embrittlement was reported at temperatures rangingfrom 300 to 400°C after long term aging. Since thesematerials have to serve for a very long period of timeduring the life span of the plants, understanding theembrittlement mechanism at slightly above the servicetemperature is very important.

The precipitation sequence in 17-4 PH stainless steelbegins with formation of coherent copper precipitates,which occurs during the tempering treatment before ser-vice. These coherent particles were reported to trans-form to incoherent fcc-Cu particles after long term ag-ing at temperatures around 400°C [3]. In addition, sincethe Cr concentration in 17-4 PH is within the spinodalline, phase decomposition of the martensite into the Fe-rich α and the Cr-enriched α’ is expected on aging be-low 450°C. Much work has shown that stainless steelsare embrittled when α’ phase precipitates by spinodaldecomposition [11]. Such α’ embrittlement is anticipatedin the 17-4 PH stainless steel as well.

Several studies on the effect of aging 17-4 PH stain-less steel were carried out [3-6]. Early work by Anthony[4] proposed mechanical properties of 17-4 PH are in-fluenced by precipitation of α’ phase, but no direct evi-dence for α’ precipitation was presented. Later, Jackand Kalish [3] observed copper precipitation on agingand correlated it to mechanical property changes; how-ever, there was no mention of phase decomposition inthe martensite phase. More recently, Yrieix andGuttmann [10] reported that 17-4 PH stainless steel ex-hibits high susceptibility to aging embrittlement at400°C, and they concluded that it was essentially dueto α’ precipitation. In their study, however, no micro-structural observation results were shown. Employingatom probe field ion microscopy (APFIM) and trans-mission electron microscopy (TEM), Miller and Burke[6] showed direct evidence for a’ precipitation after ag-ing at 482°C. They also reported that significantamounts of iron, nickel and manganese were containedin the ε-Cu precipitates even in the overaged condition.However, their aging temperature is rather high com-pared to the service condition of 17-4 PH steel.

This study aimed to carry out a more complete char-acterization of microstructures in 17-4 PH stainless steelat various stages of heat treatment, i.e., after solutionheat-treatment, tempering at 580°C for four hours, andlong term aging at 400°C, in order to obtain a betterunderstanding of the embrittlement phenomena on ag-ing.

II. EXPERIMENTAL PROCEDURES

The chemical composition of the alloy used in thisstudy 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-treatedat 1050°C for 1 h and subsequently water quenched.The solution treated samples were then aged at 580°Cfor 4 h (tempering). This heat treatment causes precipi-tation of coherent Cu precipitates in the martensite phase

Microstructural Evolution in a 17-4 PH Stainless Steelafter 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 heat-treatment, tempering at 580°C and long term aging at 400°C, have been studied by atom probe field ionmicroscopy (APFIM) and transmission electron microscopy (TEM). The solution treated specimen con-sists largely of martensite with a small fraction of δ-ferrite. No precipitates are present in the martensitephase, while spherical fcc-Cu particles are present in the d-ferrite. After tempering for 4 h at 580°C, coher-ent Cu particles precipitate in the martensite phase. At this stage, the Cr concentration in the martensitephase is still uniform. After 5000 h aging at 400°C, the martensite spinodaly decomposes into Fe-rich αand 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 spinodaldecomposition of the martensite phase, G-phase precipitation occurs after long-term aging.

M. MURAYAMA, Researcher, and K. HONO, Head of 3rd Labo-ratory, are with Materials Physics Division, National Research In-stitute for Metals, Tsukuba 305-0047, Japan. Y. KATAYAMA, iswith Heavy Apparatus Engineering Laboratory, Toshiba Corpora-tion, Yokohama 230-0045, Japan

Manuscript submitted April 21, 1998.


tempering treatment, the specimen was aged for 100and 5000 hours at 400°C.

For atom probe analyses, a locally built reflectron-type energy compensated time-of-flight atom probe(1DAP) and a three-dimensional atom probe (3DAP)equipped with CAMECA’s tomographic atom probe(TAP) detection system [12] were used. One disadvan-tage of the 3DAP was its poor mass resolution, becauseit was not equipped with an energy compensator forimproving mass resolution. The mass resolution of the3DAP used in this study was limited to m/∆m~200 fullwidth at 10 pct maximum (FW10 pct M), which is sig-nificantly lower than that obtained using an energy com-pensated atom probe (~500 FW10 pct M). Thus, Fe2+

and Mn2+ ions, which have similar mass-to-charge ra-tios were not distinguished in the 3DAP analyses. Thus,detailed spatial information provided by 3DAP analy-sis was complemented by 1DAP analysis with a highmass resolution. Field ion microscopy images wereobserved at temperatures of 30 - 60 K with Ne as animaging gas, and atom probe analyses were carried outat a specimen temperature of about 30 K, under a UHV(~1x10-10 torr) condition, with a pulse fraction (V

p/V

dc)

of 20 % and a pulse repetition rate of 600 Hz. Micro-structures of the specimens were examined with aPhilips CM200 transmission electron microscope(TEM), operated at 200 kV. High resolution transmis-sion electron microscope (HRTEM) observations werecarried out using a JEOL JEM-2000EX, operated at 200kV. Thin foils for TEM were prepared by grinding theslices to a thickness of about 100 mm, then by twin-jetelectropolishing using a 5 pct perchloric acid-acetic acidsolution at 287 K. For long-term aged specimen, ionbean thinning was employed for thin foil preparation,because it was found that Cu particles are preferentiallydissolved by electropolishing.

III. RESULTS

A. Mechanical properties

Figure 1 shows the influence of aging times on yieldstrength at 350 and 400°C. For both temperatures, anincrease in yield strength occurs after 10 hours aging,and the strengthening response is much faster at 400°C.Increase in the yield strength is almost saturated after10,000 hours aging, and 80 and 90 pct of strengtheningis achieved after 100 and 1000 hours aging respectively.Values of yield strength, tensile strength, elongation andCharpy V-notch energy absorption measured before andafter 5000 h aging at 400°C are summarized in Table I.Increases in yield strength, tensile strength occur afterlong term aging accompanied by decreases in elonga-tion and Charpy V-notch energy absorption. This indi-cates that embrittlement occurs as a result of long termaging.

Yield strength (Mpa)

Tensilestrength

(Mpa)

Elonga-Tion (%)

Charpy V-notch energy

absorp-tion (J)

Pre-aged alloy(580°C x 4h)prolonged aged alloy(400°C x 5000h)

Charpy V-notch energy absorption was measured at 0°C.

895 1085 23 107

1362 1434 8.3 3

Table I Changes in mechanical properties of 17-4 PHby aging

1600

1400

1200

1000

800

Yie

ld S

tren

gth

/ M

Pa

Tem

per

ed 10 100 1000 10000

Aging time / hr

400 °C

350 °C

Fig. 1 0.2% yield stress of 17-4 PH stainless steel as a function ofaging time at 400｡C.

NbC

δ-ferrite

1µm

Fig. 2 TEM bright field image of the martensite phase in 17-4 PHstainless steel after solution heat-treatment. The predominant phaseis lath martensite. Grains of NbC and δ-ferrite are indicated.

and provides balanced strength and toughness as shownin Table. I. This is the typical condition of 17-4 PH stain-less steel before use as a structural material. After this


B. Solution treated microstructure

A solution-treated 17-4 PH stainless steel is com-posed largely of martensite with a minor fraction of d-ferrite as shown in Figure 2. The martensite phase isconsist of a lath structure containing a very high den-sity of dislocations. There is no evidence of precipi-tates which suggests the martensite phase is supersatu-rated with Cu and Cr in the solution-treated condition.

On the other hand, a high density of fine precipi-tates are observed in the d-ferrite phase as shown inFigure 3 (a). In the bright-field image, the precipitates

are spherical and each precipitate appears to be associ-ated with dislocation. Absence of strain contrast andpresence Moire fringe indicate that the particles are in-coherent with the bcc matrix. In fact, selected area dif-fraction (SAD) pattern taken slightly inclined from the[111] zone show {020}

fcc reflections, indicating that the

precipitates are fcc-Cu. The orientation relationship(OR) between the particle and the d-ferrite does notmatch perfectly with the K-S relationship, but has aslight deviation from it. Energy dispersive X-ray (EDX)analysis results shown in Table II indicate that signifi-cantly higher Cu concentration is recorded from theprecipitate region than from the other regions. This alsosuggests that the precipitates are Cu, not NbC as previ-ously reported [3]. The EDX results also show that Crconcentration in the δ-ferrite is slightly higher than thatin the martensite phase.

Since the δ-ferrite is a minor constituent phase inthis steel, atom probe analyses were carried out onlyfrom the martensite phase. Figure 4 shows atom probeconcentration depth profiles of the martensite phase inthe solution-treated specimen. Horizontal broken linesshow average concentration, c

av, and statistical errors

expected from the number of atoms used for determin-ing local concentrations, c

av ± 2s, where s is the stan-

dard deviation. The number of atoms is linearly corre-

(a)

(b)

400nm

011

020fcc

Fig. 3 (a) TEM bright field image of the δ-ferrite phase in 17-4PH stainless steel after solution treatment. Fine Cu precipitates areobserved. Some of them are apparently associated with disloca-tions. (b) SAD pattern taken slightly inclined form the [111] zone.

　　　　 Cr Ni Si Cu Mnprecipitate 20.9 3.52 0.94 11.8 0.43δ-ferrite 22.7 4.00 1.43 3.0 0.40martensite 17.3 4.93 1.11 3.63 0.34

Table II. EDX analysis results of Cu-rich precipitate in-ferrite, -ferrite and martensite phase in the solution

treated condition (all in at. %).

100

80

60

40

20

0

30

15

0

10

0

30

15

0

30

15

00 50 100 150 200 250 300

Fe

&C

r / a

t.%

Ni

Si

Mn

Cu

Number of Atoms / x50

Fig. 4 1DAP concentration depth profiles of Fe, Cr, Ni, Si, Mnand Cu in the martensite phase in 17-4 PH stainless steel after solu-tion treatment.


lated with the depth scale of the analysis, and the totaldepth of this analysis is estimated to be approximately60 nm. This data shows that the martensite in the solu-tion heat-treated specimen is a supersaturated solid so-lution containing all solute atoms homogeneously.

C. Tempered microstructure

Figure 5 shows a bright field TEM image of themartensite phase after aging for 4 h at 580°C. Bright-field image does not give any clear contrast correspond-ing to fine coherent Cu precipitates expected in thisstage. This suggests that Cu precipitates, if any, is stillcoherent and they do not cause large strain contrast dueto the small strain field around the precipitate. In fact,the SADP taken from a martensite lath does not showany evidence for a secondary phase, suggesting thatthere is no precipitates with the distinct structure dif-ferent from the bcc matrix. Figure 6 shows 3DAP el-emental mapping of Cu and Cr obtained from the mar-tensite phase. The Cu mapping clearly shows that thereis a small spherical particle enriched with Cu. On theother hand, the Cr mapping shows that the distributionof Cr atoms is uniform martensite phase. The Cu el-

emental map clearly shows that there is Cu enrichedprecipitate in this stage. As the SADP (Figure 5(b) and(c)) does not show any evidence for presence of thesecondary phase, the Cu enriched precipitate observein the 3DAP data is believed to be fully coherent bcc-Cu. In order to quantify the concentration of the par-ticle, concentration depth profile was measured fromthe selected region near the Cu-rich precipitate as shownin Fig. 6 (c). The chemical composition of the Cu-richprecipitates has been found to be 55 at.pct Cu, 30 at.pctFe, 10 at. pct Cr, 5 at. pct Ni. It is seen that Cr and Niare rejected from the Cu-enriched particle slightly. Itshould also be noted that the concentration of Cu in theparticle is significantly lower than that expected fromthe equilibrium e-Cu.

D. Aging for 100 h at 400°C

Figure 7 (a) and (b) show 3DAP elemental mappingof Cu and Cr obtained from the martensite phase in thespecimen aged for 100 h at 400°C. The Cu mappingshows that a high density of Cu-rich precipitates of ap-proximately 3 nm in diameter is present. The Cr map-ping shows that the distribution of Cr atoms is no longeruniform and fluctuations of Cr concentration occur.Concentration depth profiles of Fe, Cr, Ni, Si and Cuwere measured from a selected region cutting two Cu-rich precipitates as shown in Figure 6(c). The concen-tration of Cu in the Cu-rich precipitates is approximately70 at. pct Cu, which is much higher than that observedin the tempered specimen, but it is still lower than theequilibrium concentration of the ε-Cu. The presence offluctuations in Cr concentration indicates that the phasedecomposition occurs in the martensite phase, and itdecomposes to Fe-rich a and Cr-enriched α’ phases. TheCr concentration in the α’ phase is only 25 at. pct andthis is significantly lower than that of equilibrium α’phase, suggesting that it is still in an initial stage of thedecomposition process. It should be noted that there isno indication of partitioning of Ni.

E. Aging for 5000 h at 400°C

Figure 8 (a) shows a dark field image excited usingthe 1/2(110) reflection and Figures 8 (b) – (d) shows[001], [011], [111] zone SAD patters, respectively ob-tained from the martensite phase aged for 5000 hoursat 400°C. Extra reflections indicate an ordered phasehaving a cube-on-cube orientation relationship precipi-tates after prolonged aging. The dark field image whichwas taken using the 1/2(110) reflection shows a highdensity of fine, ordered particles are dispersed in themartensite matrix. The diffraction pattern was found tobe consistent with the G phase (structure type D8

a, space

group Fm3m ) reported in aged duplex stainless steels[23,26,27] and type 308 stainless steels [19].

Figure 9 shows atom probe concentration depth pro-files obtained from the martensite phase in the alloy

(b)

110

110(c)

110

011

100nm

(a)

Fig. 5 (a) TEM bright field image, (b) [001] and (c) [111] SADpattern of the martensite phase tempered at 580°C for 4 h.


aged for 5000 hours at 400°C. In addition to Cu-en-riched precipitates, significant fluctuations in Cr con-centration are observed. This shows that phase decom-position occurs in the martensite phase, and the mar-tensite decomposes to Fe-rich α and Cr-enriched α’phases. The Cr concentration in the α’ phase is approxi-mately 40 at. pct, which is again significantly lowerthan that expected from thermal equilibrium (~90 at.pct). Thus, it is believed that this stage is still in themiddle of the decomposition process, and has notreached the equilibrium. The Ni concentration profiledoes not show any tendency of partitioning of Ni to thea phase, unlike previous report by Danoix in the agedferrite phase of duplex stainless steel [26]. It appears

that Cu enriched precipitates do not have any correla-tion with the Cr concentration fluctuations. From theconcentration profile, two types of Cu precipitates areobserved. One is composed of Cu only, and the other isenriched with Ni, Si and Mn, as well. The apparent Cuconcentration of the latter is ~20 at. pct Cu. This obser-vation is similar to the solute partitioning in an ultrafinecopper-enriched zone in a neutron irradiated A533Bsubmerged 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 massspectrum of Mn2+ overlaps with that of Fe2+. The Crmapping shows the concentration of Cr fluctuates. TheCr enriched regions appear to be interconnected, which

100

80

60

40

20

0

30

15

0

60

40

20

0

30

15

0

0 2 4 6 8 10

Fe

Cr

Fe

& C

r / a

t. %

Ni

Si

Cu

Depth / ~nm

(c)

(a) Cr

(b) Cu

~26nm ~26nm

~16

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 Curich precipitate is observed. (c) Concentration depth profile through a Cu precipitate.


is a typical feature of spinodal decomposition. The pe-riodicity of the fluctuation is on the order of 3 nm. TheCu mapping shows that Cu-enriched particles approxi-mately 8 nm in diameter are present. In direct contactwith one of the Cu particle, a Ni and Si enriched par-ticle is observed as indicated by an arrow in Figure10(b). Such fine particles are always observed in con-tact with Cu precipitates, and they are believed to bethe G-phase based on the TEM observation shown inFigure 8. Figure 10 (c) shows concentration depth pro-files across the Cu precipitate and the G-phase indi-

(b)

Fe

Cr

100

80

60

40

20

0

30

15

0

70

50

30

100

30

15

0

0 5 10 15 20 25

Depth / ~nm

Fe

& C

r / a

t. %

Ni

Si

Cu

(a)

Cr

Cu ~30nm

~9n

m

Fig. 7 3DAP elemental mapping of the martensite phase aged at580°C for 4 h. (a) The phase decomposition into Cr enriched anddepleted region occur in the martensite phase. (b) Fine sphericalCu rich precipitate. (c) Concentration depth profile obtained fromthe selected region near the Cu precipitate.

100nm

(a)

(d)

011

110

(b)

110

110

(c) 011

200

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. Thedark field image was taken using the 1/2(011) reflection.

Number of Atoms / x50

100

80

60

40

20

0

Fe

&C

r / a

t.%

30

15

0

30

15

0

030

15

0

10

5000 1000 1500

Ni

Si

Mn

Cu

Fig. 9 1DAP concentration depth profile of the martensite phaseaged at 400°C for 5000 h. Ni, Si, and Mn appear to be partitionedinto the Cu precipitate.


cated by the arrow in Figure 10 (b). The Cu precipitatecontains approximately 95 at. pct Cu, which is close tothe equilibrium concentration of ε-Cu. The composi-tion of the G-phase is approximately 55 at. pct Ni, 25

at. pct Si, 20 at. pct Fe. In addition, Mn is also enrichedin the G-phase based on the 1DAP result shown in Fig-ure 9 (Mn ions can be differentiated from Fe atoms us-ing 1DAP). According to the concentration depth pro-

(b) Cu, Ni and Si

Ni + Si Cu

~15nm~15nm

~18

nm

(a) Cr

(c)

100

80

60

40

20

0

50

0

Fe

& C

r / a

t. %

Ni

Si

Cu

0 2 4 6 8 10

Depth / ~nm

100

80

60

40

20

0

100

80

60

40

20

0

12 14

G phaseCu precipitate

Fe

Cr

Fig. 10 3DAP elemental mapping of the martensite phase aged at 400°C for 5000 h. (a) The phase decomposition into Cr enriched anddepleted region occur in the martensite phase. (b) Large dots correspond to Cu atoms and small dots to Ni and Si atoms. (c) Concentrationdepth profile across the Cu precipitate and the Ni-Si enriched particle obtained from the selected region near the Cu precipitate (indicated byarrow).


file obtained by the conventional atom probe (Figure9), Ni, Si and Mn atoms appear to be partitioned to theCu enriched particle indicated by the dashed lines onthe left side of the figure. However, this result is prob-ably an artifact caused by the convoluting effect of theprobe hole covering both Cu and G-particles becauseNi, Si and Mn enrichment is not observed at the otherCu enriched region in Figure 9. In the latter case, theprobe hole covered only the Cu particle and missed theG-phase which was adjacent to the Cu particle. Theseresults demonstrates that employment of the 3DAP tech-nique provides more accurate data for characterizingthe morphological feature of fine precipitates embed-ded in the matrix.

Figure 11 (a) shows an HREM image taken alongthe [111] zone axis of the martensite phase in the alloyaged at 400°C for 5000 hours. In this image, fringe con-trast having a periodicity of two (011)

bcc planes is ob-

served. This is consistent with the fringe contrast ex-pected from the G-phase (Ni

16X

6Si

7, X=Fe, Mn,

Si, Fm3m , a=0.406 nm). Furthermore, a small particleis observed adjacent to the G-phase, which is believedto be a Cu precipitate. A microdiffraction pattern takenfrom the region with the Moire fringes is shown in Fig-ure 11 (b). The [111] microdiffraction pattern shows

the extra reflections near the {110}bcc

reflections. Theδ-spacing calculated from the extra reflections is ~0.179nm which corresponds to the spacing of the {020}

fccplanes in fcc-Cu (0.180

nm). The orientation relation-

ship (OR) between the particle and the martensite al-most matches with the K-S relationship.

IV. DISCUSSION

This study has clarified evolution of microstructurein a 17-4 PH stainless steel during long term aging at400°C. Emphasis is given to characterization of chemi-cal features of the precipitates which appear in the mar-tensite phase after prolonged aging.

The ε-Cu precipitates have been found in the δ-fer-rite after the solution heat treatment. The appearanceof these precipitates in the δ-ferrite was unaffected bysubsequent aging. Rack and Kalish [3] reported similarmicrostructural feature in δ-ferrite, but they attributedthem to NbC precipitates. In this study, NbC were ob-served at martensite lath boundaries with much largersize as shown in Figure 1, and we believe that the fineprecipitate in the δ-ferrite observed in the previous studyas well as in this study are the ε-Cu. Precipitation of Cuin the Fe-Cu binary system has been a subject of nu-merous studies [14-18], and it is well established that theinitial Cu-enriched precipitate is perfectly coherent withthe bcc matrix, while large overaged precipitates havean fcc structure with the K-S OR. However, binary al-loys were all solution heat-treated in the single phaseferrite region around 900°C, while the 17-4 PH stain-less steel is solution heat-treated at much higher tem-perature around 1050°C. Diffusivity of Cu in the d-fer-rite at this temperature (D

Cu ~ 1 x 10-9 cm2/s) is more

than three orders of magnitude higher than that in theaustenite (D

Cu ~ 4 x 10-12 cm2/s), thus we believe that

Cu precipitated out from the δ-ferrite during coolingafter the solution heat treatment. Many Cu particles wereobserved along dislocations, and this would have easethe strain contrast for precipitation of the fcc ε-Cu. Onthe other hand, no indication of presence of precipi-tates are recognized in as-quenched martensite. Thesolubility of Cu in the austenite phase is up to 7 at. pctat 1100°C. However, the diffusivity of Cu in the auste-nite is orders of magnitude small than that in the ferritephase. Thus, the precipitation kinetics are much slowerin the austenite phase and Cu can be quenched in themartensite phase from the solution heat-treatment tem-perature. After tempering at 580°C for 4 hours, fine Cu-rich precipitates were detected in the martensite phaseusing 3DAP. The Cu content in the Cu-rich precipitatesis significantly lower than the equilibrium value for ε-Cu. They reported that the average copper concentra-tion of small precipitates is approximately 50 at. pct inthe earliest stage of the precipitation. The copper con-centration in the fine, Cu-rich particles which precipi-tate during tempering is in good agreement with theearly study by Goodman et al. [16]. Unlike in the d-fer-

(a)

G phase Cu precipitate

2nm

(b)020fcc

011

Fig. 11 (a) HREM image taken at the [111] zone axis of the mar-tensite phase in the alloy aged at 400°C for 5000 h. (b) micro-dif-fraction pattern obtained from the region with the Moire fringe.


rite, these Cu-rich precipitates observed in the marten-site phase after tempering are bcc-Cu.

The martensite phase decomposes into Fe-rich α andCr-enriched α′ phases by the spinodal mechanism. Inaddition, fine precipitates of the G-phase has been foundin the martensite phase in direct contact with the Cuprecipitates after 5000hours aging at 400°C. Concen-tration of Cr in the α′ phase is approximately 25 at. pctafter 100 hours, then reach 40 at. pct after 5000 hoursaging. This concentration is significantly lower than thatexpected from the binodal line and is thought to be stillin the decomposition process. Since no precipitation ofthe G-phase was observed after 100 hours aging, theincrease in yield strength up to 100 hours is attributedto the spinodal decomposition rather than the forma-tion of the G-phase. G-phase precipitation was reportedin type 308 stainless steels [9, 19] and maraging steels [20]

as a grain boundary phase. In recent years, it was shownthat the G-phase exists in various Fe-Cr-Ni alloys as adispersed phase in the interior of the grains [21-27]. Forexample, Auger et al. [27] reported phase separation andprecipitation of the G-phase in the ferrite phase of du-plex stainless steel. They concluded that the nucleationof the G-phase takes place at the ‘interface’ betweenCr-rich a¢ and Fe-rich a. However, in the case of the17-4 PH stainless steel, we have found that G-phaseprecipitation occurs in intimate contact with ε-Cu pre-cipitates after the decomposition of the martensite haveprogressed. This result indicates that the Cu precipi-tates provides heterogeneous nucleation sites for G-phase formation. One reason for this is that the Cu/martensite interface is a more suitable site for nucle-ation of the G-phase than the α/α’ interface. The darkfield image of the long-term aged martensite phaseshows the G-phase is dispersed uniformly in the grainrather than grain boundary as commonly observed inother stainless steels. This suggests that the G-phaseprecipitation itself does not contribute to theembrittlement. In fact, 80 pct of total yield strength in-crease occurs before precipitation of the G-phase isobserved (100 hours aging at 400°C).

Our atom probe results have shown that the atomicratio of Ni to Si in the G-phase is 6:4. In addition, Mnand Fe are also contained in the G-phase. The ternarysilicide designated as the G-phase is known to be fcchaving 116 atoms in a unit cell. The lattice parameter is~1.12 nm. The structure of this phase is isotypic withTh

6Mn

23 (structure type D8a, space group Fm3m), and

the ideal composition was proposed as X6Ni

16Si

7, where

X is typically Ti, but can be substituted for other transi-tion element. Considering the qualitative nature of theatom probe result, the determined composition (55 at.pct Ni, 25 at. pct Si, 20 at. pct Fe and some Mn) isreasonably close to the ideal stoichiometry of X

6Ni

16Si

7when one allows for substitution of Mn and Fe forX [28].

In this study, partitioning of Ni in the α phase after

spinodal decomposition of the martensite was not ob-served. This result is in contrast to the APFIM resultsby Danoix et al. [26] which reported that nickel is re-jected from Cr-enriched α’ phase and partitioned intothe Fe-rich α phase in the aged ferrite phase of duplexstainless steel. In the equilibrium condition, the solu-bility limit of Ni in α and α’ were estimated to be 4.1and 0.01 at. pct, respectively, by the ThermoCalc soft-ware. Thus, Ni should partition into the α’phase in 17-4 PH as well. This suggests Ni partitioning does notoccur until decomposition progresses further and thedriving force for the partitioning reaction increases.Similar delayed partitioning behavior was reported forAl in an Fe-Cr based alloy [29]. As the Ni concentrationof the duplex stainless steel was higher than that in 17-4 PH, the driving force for Ni partitioning is much higherin the duplex stainless steel. Thus the absence of con-current partitioning of Ni with spinodal decompositionin 17-4 PH is reasonable.

V. CONCLUSIONS

The microstructural features of 17-4 PH stainlesssteel at various stages of heat-treatment has been in-vestigated by APFIM and TEM. The main results areas follows:

1. The fcc Cu-rich particles precipitate in the δ-ferritegrains during cooling after the solution heat treat-ment. The precipitates have the K-S orientation re-lationship with the δ-ferrite matrix.

2. Tempering at 580°C for 4 hours results in precipita-tion of coherent bcc particles in the martensite phase.The composition of the precipitates is approximately60 at. pct Cu and Cr, Ni and Si are rejected from thebcc-Cu. The Cr concentration is homogeneous inthe martensite phase at this stage.

3. After 100 hours aging at 400°C, evidence forspinodal decomposition of the martensite phase intoFe-rich a and Cr-enriched a’ phase is found. The Crconcentration in the a’ phase is approximately 25 at.pct, significantly lower than the equilibrium value.The concentration of Cu in the Cu-enriched precipi-tate is approximately 70 at. pct at this stage.

4. Spinodal decomposition of the martensite phaseprogresses further after 5000 hours aging. The Crconcentration of the α’ phase is 40 at. pct at thisstage. The structure of the coarsened Cu precipitatesis fcc and their concentration is almost 100 at. pctCu. Fine particles of G-phase containing approxi-mately 60 at. pct Ni, 25 at. pct Si and some Fe andMn are present in direct contact with the Cu pre-cipitates, suggesting that G-phase is heterogeneouslynucleated at the martensite/Cu interface.

5. The increase in hardness and yield strength of the17-4 PH stainless steel after aging at 400°C is mostlycaused by spinodal decomposition of the martensitephase. G-phase precipitation does not appear to make


a significant contribution to the embrittlement dur-ing aging.

ACKNOWLEDGMENTS

The authors thank Professor W.T. Reynolds, Jr., Vir-ginia Polytechnic Institute and State University, forvaluable discussions. This work was supported by theFrontier Research Center for Structural Materials,NRIM.

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