Effect of Cu on the Precipitation of Deleterious Phases and the MechanicalProperties of 27Cr­7Ni Hyper Duplex Stainless Steels

Soon-Hyeok Jeon, Il-Jeong Park, Hye-Jin Kim, Soon-Tae Kim,Young-Kook Lee and Yong-Soo Park+

Department of Material Science and Engineering, Yonsei University,134 Shinchon-dong, Seodaemun-gu, Seoul 120-749, Republic of Korea

Cu addition to the base alloy reduces the total amount of deleterious phases such as chromium nitride and sigma and chi phases. Inparticular, Cu addition to the base alloy results in pronounced suppression of the amount of sigma phase whereas it slightly facilitates theprecipitation of chromium nitride and chi phase along the phase boundaries and within ferrite grains. During the initial stage of aging, thepreferential precipitation of chromium nitride and chi phase seems to be closely associated with the retardation of the precipitation of the sigmaphase. The preferential precipitation of chromium nitride and the chi phase inhibits the nucleation and growth of the sigma phase by depletingthe Cr adjacent to the chromium nitride particles and depleting the Mo and Wadjacent to the chi phase. Thus, the addition of Cu to the base alloyreduces its embrittlement owing to the delayed precipitation of these deleterious phases. [doi:10.2320/matertrans.M2013471]

(Received December 26, 2013; Accepted March 26, 2014; Published May 2, 2014)

Keywords: copper, duplex stainless steel, sigma phase, chi phase, chromium nitride

1. Introduction

Super duplex stainless steels (SDSSs) are increasinglybeing used in various applications, including in power plants,desalination facilities, and chemical plants. This is owingto their high resistance to pitting and crevice corrosion,excellent mechanical properties, and relatively low cost, incontrast to other high performance materials such as superaustenite stainless steels.1­3)

However, when used in shell and tube exchangers, SDSSsexhibit corrosion resistance that is insufficient to allow forhigh temperature operations and a long service life. Hence,hyper duplex stainless steels (HDSSs), which exhibit highresistance to pitting corrosion, combined with improvedmechanical properties, have been developed.

When duplex stainless steels (DSSs) are aged at 600­950°C, deleterious phases such as chromium nitride, as wellas secondary (i.e., sigma (·) and chi (»)) phases, tend toprecipitate in them.4­9) This precipitation of secondary phasesand chromium nitride leads to a significant reduction in thecorrosion resistance of the steels as well as a deterioration oftheir mechanical properties owing to their inherent brittlenessand the formation of Cr-, Mo-, and W-depleted regionsaround the phases.10­15)

DSSs contain a significant amount of Cr, Mo, and Wwhich improves corrosion resistance. These alloying ele-ments facilitate precipitation of secondary phases such asthe · and » phases. In addition, for DSSs, the tendency tosecondary phases precipitation is crucial since the existenceof the ferrite (¡) phase will enhance the kinetics forprecipitation of secondary phases. The secondary phasespreferentially precipitate into the ¡ phase due to the higherCr and Mo concentration in the ¡ phase.16) A fundamentalreason why the secondary phases preferentially grow into the¡ phase is that the ¡ phase is thermodynamically meta-stableat temperature where the secondary phases precipitate.17,18) Inaddition, the diffusion rates of the alloying elements in the ¡

phase are 100 times faster than the corresponding values inthe austenite (£) phase.19,20)

The effects of various alloying elements on the precip-itation of secondary phases such as · and » phases have beeninvestigated previously. The addition of C and N in ferriticstainless steels (FSSs) retards the formation of the · phase byincreasing the incubation period.21) The addition of N reducesthe tendency of · phase formation in the ¡ phase since itresults in the removal of Cr from the solution through theformation of chromium nitride.22) Kim et al.23) reported thatthe addition of a small amount of Ce (55­110 ppm) to HDSSsresults in the homogeneous distribution of Ce in the alloymatrix and delays the precipitation of secondary phases byreducing the diffusion rates of Cr, Mo, and W. Park et al.24)

showed that partially substituting W for Mo retarded theformation of the · phase in FSSs. This was due to an increasein the tendency of formation of a » phase with a highernucleation efficiency and lower growth rate than those of the· phase. Nana and Cortie25) reported that the addition of Cuin FSSs suppresses the precipitation of the · phase since theexcessive Cu is shifted away from the reaction ¡/· interface.Smuk et al.26) showed that the addition of Cu results in Cuprecipitates in the ¡ phase during slow cooling and retardsthe precipitation of the · phase; this is due to the fact thatthe Cu precipitates pin the moving · phase boundaries inDSSs.

However, the effects of Cu addition on the precipitation ofchromium nitride and the · and » phases as well as on themechanical properties of HDSSs have not yet been inves-tigated. In particular, the mechanisms by which the additionof Cu affects the precipitation of chromium nitride and the ·and » phases in HDSSs during the initial stage of aging havenot been elucidated.

Therefore, in this study, we employed focused ion beam(FIB) milling and transmission electron microscopy (TEM) tostudy the microstructural changes induced in HDSSs duringthe initial aging stage by the addition of Cu. The FIB millingtechnique was used to fabricate the TEM samples, allowingus to prepare samples from specific sites of interest such that+Corresponding author, E-mail: corrus@yonsei.ac.kr

Materials Transactions, Vol. 55, No. 6 (2014) pp. 971 to 977©2014 The Japan Institute of Metals and Materials EXPRESS REGULAR ARTICLE

the precipitation behavior of chromium nitride and · and» phase could be investigated in depth. This is permittedfor greater flexibility in the TEM-based measurements. Inaddition, to elucidate the effects of Cu addition on theprecipitation of chromium nitride and the · and » phasesas well as the mechanical properties of HDSSs, X-raydiffraction (XRD) analysis, scanning electron microscope-energy dispersive X-ray spectroscopy (SEM-EDS) analyses,and TEM-EDS, image analysis, tensile tests were alsoperformed.

2. Experimental Procedures

Ingots weighing 50 kg with dimensions 150 by 150 by300mm (width by length by height) were manufacturedusing a high frequency vacuum induction furnace. After theseingots were hot rolled in the range of 1333 to 1523K, platesof 6mm thickness were manufactured. The chemicalcompositions of the alloys are presented in Table 1. Theexperimental alloys were cut and solution heat-treated for5min per 1mm thickness at 1363K and then quenched inwater. The specimens were then isothermally aged at 1123Kfor 10 and 30min.

To observe the microstructures of the alloys, they wereground to 2000 grit using SiC abrasive papers, polishedsurface using diamond paste. The sample was ultrasoniccleaned in acetone and distilled water to remove anyimpurities from the polished surface of the sample. The ·

and » phases and chromium nitrides were observed usingSEM in backscattered electron mode (BSE). In addition, thechemical compositions of the secondary phases and chro-mium nitrides were analyzed by an EDS attached to the SEM.XRD analysis was performed on specimens containing arelatively large fraction of precipitates for the phaseidentification. FIB milling technique for the preparation ofTEM specimens was conducted. The » phase and chromiumnitrides were analyzed using a TEM. The line profiles of theCr, Mo and W in the » phase and chromium nitrides weremeasured using an EDS attached to the TEM.

Tensile tests were conducted at room temperature toinvestigate the effects of secondary phases on the ductility ofthe alloys under constant strain rate of 0.01/s. Tensilespecimens were cut parallel to the rolling direction of thesheets 3mm thick with a gage section 25mm long and6.35mm wide.

3. Results and Discussion

3.1 MicrostructuresFigure 1 shows the microstructures of solution-annealed

alloy specimens and aged alloy specimens at 1123K for 10and 30min. The solution-annealed alloy specimens had £

phase and ¡ phase without secondary phases. The £ phasecan be seen as an isolated phase on the background of the¡ phase, which looks relatively dark (Figs. 1(a) and 1(d)).Meanwhile, after the specimens had being aged at 1123K for10min, it was found that precipitates were formed in thespecimens and that these precipitates formed continuousnetworks along the grain boundaries while also appearingrandomly within the grains (Figs. 1(b) and 1(e)). Theseprecipitates were revealed the » and · phases and chemicalcompositions of these were confirmed by SEM-EDS analy-ses, as shown in Table 2. Only precipitates greater than 1 µmin size were selected in order to minimize the influence of thematrix on the analysis results. A few precipitates of the »

phases formed continuously along the phase boundariesbetween the £ and ¡ phases, and the · phase appearedrandomly within the grain of the ¡ phase, in keeping with theeutectoid reaction. In the case of the specimens that wereaged at 1123K for 30min (Figs. 1(c) and 1(f )), the total size

Table 1 Chemical compositions of the experimental alloys (mass%).

Alloy C Cr Ni Mo W Si Mn Cu S N Fe

Base 0.020 27.01 7.00 2.52 3.28 0.35 0.88 ® 0.003 0.35 Bal.

1.5Cu 0.017 26.91 6.59 2.50 3.30 0.33 0.94 1.45 0.004 0.38 Bal.

35 μm

σχ

σχ

σ

χ

σ

χ

35 μm

35 μm 35 μm

Base alloy-10 min. Base alloy-30 min.

1.5Cu alloy-10 min. 1.5Cu alloy-30 min.

(c)

(e) (f)

35 μm1.5Cu alloy-solution annealed

(b)

35 μmBase alloy-solution annealed

(a)

α

γ

α

γ

(d)

Fig. 1 SEM-BSE images of the alloys aged at 1123K for 10 and 30min: (a) solution-annealed base alloy, (b) base alloy aged at 1123Kfor 10min, (c) base alloy aged at 1123K for 30min, (d) solution-annealed 1.5Cu alloy, (e) 1.5Cu alloy aged at 1123K for 10min and(f ) 1.5Cu alloy aged at 1123K for 30min.

S.-H. Jeon et al.972

of the secondary (» + ·) phases at the grain boundaries aswell as that within the ¡ grains increased. Although thesecondary (» + ·) phases grew further in both the exper-imental alloys, the precipitation rate and type of secondaryphases in the case of the base alloy were different from that inthe case of the 1.5Cu alloy. In particular, in the case of thebase alloy, in numerous instance, the · phase formed withinthe ¡ grains and the » phase precipitated continuously alongthe boundaries between the £ and ¡ phases. However, in thecase of the 1.5Cu alloy, the · phases did not grow to asignificant degree at the grain boundaries and within the ¡

grains, but, in a number of instances, the » phase grew furtheralong the phase boundaries between £ and ¡ phases andwithin the ¡ grains. In a previous study, it was found that theaddition of Cu to the HDSSs facilitates the precipitation rateof » phase owing to the increase in the activity of W andretards the precipitation of · phase owing to the decrease inthe activity of Mo.27)

Figure 2 shows the XRD analysis of the alloys aged at1123K for 30min. With increasing aging time, the peak of¡ phase is led to become lower conspicuously, graduallyderiving · phase. After aging at 1123K for 30min, the peakof · phase of the base alloy is higher than that of 1.5Cu alloy.On the other hand, the peak of » phase of the base alloy islower than that of 1.5Cu alloy. Small amounts of · and »

phases are not detected by XRD analysis due to overlapbetween secondary phases with ¡ phase and, mainly, £ phasereflections. Nevertheless, the · phase reflections such as(112), (212), (411) and (331) are led to become higherconspicuously.

Figure 3 shows SEM-BSE images of chromium nitrideparticles in the alloy specimens aged at 1123K for 10 and30min. In the case of the base alloy, in numerous instances,the · phase formed within the ¡ grains and the » phaseprecipitated continuously along the boundaries between the £and ¡ phases. In addition, a few rounded and acicular nitrideparticles precipitated continuously along the phase bounda-ries (Figs. 3(a) and 3(b)). The results of the SEM-EDSanalyses indicated that these rounded and acicular nitrideparticles were of chromium nitride; this was owing to the facttheir Cr content was much higher than that of the matrix. Thechemical composition of the chromium nitride particles asdetermined by the SEM-EDS analyses is listed in Table 2.The precipitation rate and type of the secondary (» + ·)phases and chromium nitride particles in the base alloy,however, were different from those in the case of the 1.5Cualloy. In the case of the 1.5Cu alloy, the · phase did not growto a significant degree at the grain boundaries or within the ¡grains; instead, in a number of instance, the » phase andchromium nitride particles precipitated along the boundariesbetween the £ and ¡ phases and within the ¡ grains (Figs. 3(c)and 3(d)). It is known that chromium nitride particles nucleateat ¡/£ phase boundaries and within ¡ grains as well as at ¡grain boundaries (¡/¡) in the ¡ phase.28­30) Further, the ·

phase did not form or grow at the phase boundaries where thepreferential precipitation of chromium nitride and the » phaseoccurred. This result confirmed that the preferential precip-itation of chromium nitride and the » phase in theexperimental alloys inhibited the nucleation and growth ofthe · phase. The addition of Cu to the base alloy significantlysuppresses the extent of the · phase; this is due to the numberof precipitated chromium nitride particles and the » phase. Ina previous study, it was found that the addition of Cu to theHDSSs facilitates the precipitation of chromium nitride andstabilizes the alloy at high temperatures owing to the increasein the activity of Cr.31)

Figure 4 shows a TEM bright-field image, SAED pattern,and the result of the line analysis of the » phase in the 1.5Cualloy specimens aged at 1123K for 10min. The results ofthe TEM analysis demonstrated that the precipitates were »

phases in the 1.5Cu alloy specimens aged at 1123K for10min (Fig. 4(a)). Based on the SAED pattern (Fig. 4(b)),the » phase has a body centered cubic (BCC) structure. Moand W-depleted zones adjacent to the » phase were observed(Fig. 4(c)). The results show that » phases leads to Mo andW-depleted areas at the interfaces between the matrix and the» phases.

Figure 5 shows a TEM bright-field image, Selected AreaElectron Diffraction (SAED) pattern, and the result of theline analysis of the chromium nitride in the 1.5Cu alloyspecimens aged at 1123K for 10min. The results of the TEManalysis demonstrate that the precipitates were chromiumnitride particles in the 1.5Cu alloy specimens aged at1123K for 10min (Fig. 5(a)). Based on the SAED pattern(Fig. 5(b)), we realized that the chromium nitride had ahexagonal close packed (HCP) structure. A Cr-depleted zoneadjacent to the chromium nitride was observed (Fig. 5(c)).The results indicate that the formation of chromium nitrideparticles leads to Cr-depleted areas at the interfaces betweenthe matrix and the chromium nitride particles.

Table 2 Chemical compositions of the » and · phase and chromium nitrideformed in experimental alloys after aging at 1123K (mass%).

PhaseBase 1.5Cu

Cr Mo W N Cr Mo W N

Matrix 27.4 2.6 3.4 0.3 27.4 2.6 3.4 0.4

» 25.7 11.9 13.3 ® 25.4 9.8 13.9 ®

· 29.4 4.3 4.9 ® 29.5 4.2 4.9 ®

Chromiumnitride

83.5 2.9 6.3 4.4 81.1 3.42 7.2 5.1

γ(200)

α(110)γ

(111)

σ(112)

σ(411)

σ(212)

σ(331)

Χ(332)

40 42 44 46 48 50 520

100

200

300

400

500

Inte

nsi

tiy

(cp

s)

2θ (°)

Base 1.5Cu

Fig. 2 XRD analysis of the alloys aged at 1123K for 30min.

Effect of Cu on the Precipitation of Deleterious Phases and the Mechanical Properties of 27Cr­7Ni Hyper Duplex Stainless Steels 973

The preferential precipitation of chromium nitride and the» phase seems to be closely associated with the retardation ofthe precipitation of the · phase. As was found during theprevious metallographic observations, the preferential pre-cipitation of chromium nitride and the » phase occurred alongthe phase boundaries and within the ¡ grains. In particular,the » phase develops from the · phase, which acts as aprecursor. Thus, the formation of the · phase is significantly

influenced by the chemical compositions of the » phase andchromium nitride. Since the formation of the · phase in theexperimental alloys requires high concentrations of Cr, Moand W, the preferential precipitation of chromium nitride andthe » phase in the alloys during the initial stage of aginginhibits the nucleation and growth of the · phase owing to thedepletion of the Cr adjacent to the chromium nitride particlesas well as the depletion of the Mo and W adjacent to the »

(c)

0

2

4

6

8

10

12

14

16

18

20

Co

nce

ntr

atio

n o

f W

(m

ass

%)

Distance, D /nm

W depleted zone

-100 -75 -50 -25 0 25 50 75 100-100 -75 -50 -25 0 25 50 75 1000

2

4

6

8

10

12

14

16

18

20

Co

nce

ntr

atio

n o

f M

o (

mas

s %

)

Distance, D /nm

Mo depleted zone

100 nm

(a)

(113)

(110)

P113

χ (BCC)

(b)χ

Fig. 4 TEM analyses of chi phase in the 1.5Cu alloy aged at 1123K for 10min: (a) the bright field images, (b) Selected Area ElectronDiffraction (SAED) pattern and (c) the line analysis of Mo and W adjacent to chi phase.

Cr2N

σ

χ

Cr2Nσ

χ

5 μm

5 μm

Cr2N

σ χ

5 μm

Cr2N

σ

χ

5 μm

(a) (b)

(c) (d)

Base alloy-10 min. Base alloy-30 min.

1.5Cu alloy-10 min. 1.5Cu alloy-30 min.

Fig. 3 SEM-BSE images of the chromium nitrides of the alloys aged at 1123K for 10 and 30min: (a) base alloy aged at 1123K for10min, (b) base alloy aged at 1123K for 30min, (c) 1.5Cu alloy aged at 1123K for 10min and (d) 1.5Cu alloy aged at 1123K for30min.

S.-H. Jeon et al.974

phase. In the initial stage of chromium nitride and the » phaseformation, the · phase formed is small in size, because theamount of Cr, Mo, and W present for · phase formation areinsufficient around the chromium nitride particle, which havea very high Cr content, as well as around the » phase, whichhas very high W and Mo contents. The addition of Cu to thebase alloy significantly suppresses the extent of the · phase,owing to the number of precipitated chromium nitrideparticles and the » phase. Park et al.24) have reported thatthe preferential precipitation of the » phase in the FSSsduring the initial stage of aging can inhibit the nucleation andgrowth of the · phase, owing to depletion of the Mo and Wadjacent to the » phase.

Figure 6 shows the effects of Cu addition and aging at1123K on the area fraction of the » phase, the · phase andthe total secondary (» + ·) phases in the experimental alloysby computer based image analyzing system. Ten randomlydispersed BSE images were used for the quantitativemeasurement of the phase fraction. In BSE image, thesecondary phases embedded in the ¡ phase had a high Cr, Moand W content, and produced bright shadeds, such as lightgrey (·) and white (»), respectively. With an increase in theaging time, the area fraction of the » phase precipitated inthe 1.5Cu alloy becomes much higher than that in the basealloy (Fig. 6(a)). However, the area fraction of the · phaseprecipitated in the base alloy is higher than that in the 1.5Cualloy (Fig. 6(b)). Figure 6(c) shows that the area fraction ofthe total secondary phases (» + ·) precipitated in the basealloy was much higher than that in the 1.5Cu alloy. Insummary, as the aging time is increased, the addition of Cu

to the base alloy reduces the extent of the secondary phases(» + ·). In particular, the addition of Cu to the base alloysignificantly suppresses the extent of the · phase whereasit slightly increases that of the » phase. This result suggeststhat adding Cu to the base alloy effectively retards theprecipitation of the deleterious phases.

3.2 Effect of Cu addition on mechanical propertiesTensile tests were performed to measure the mechanical

properties of the alloy specimens in order to elucidate theeffect of Cu addition and the different aging times onthe precipitation of deleterious phases and the associatedmechanical properties of HDSSs. In previous studies, thesetests have shown that the embrittlement caused by theaddition of Cu may be due to the precipitation of the · and »

phases.32,33)

To investigate the effect of aging on the tensile properties,tensile tests were performed using base and 1.5Cu alloy agedat 1123K for 10min at room temperature. The tensileproperties data of base and 1.5Cu alloy aged at 1123K for30min could not be determined due to brittle failure in elasticrange. The total elongation of base alloy was greatly reducedafter aging from 29.2 to 2.5%, and that of Cu alloy wasreduced by aging from 27.8 to 10%. It is thought that theembrittlement of the alloys was apparent during plasticdeformation owing to the precipitation of secondary phases(» + ·) and chromium nitride particles. The decrease in thetotal elongation of the experimental alloys seemed to dependon the total area fraction of the deleterious phases, as shownin Fig. 7. It is realized that deleterious phases, such as » and

-300 -200 -100 0 100 200 30010

20

30

40

50

60

70

80

90

100

Co

nce

ntr

atio

n o

f C

r (m

ass

%)

Distance, D /nm

Cr depleted zone

Cr depleted zone

P0010

(0001)(0001)

(1120)(1121)

(1121)

(0000)

100 nm

(a) (b) (c)

Cr2N (HCP)

Cr2N

Fig. 5 TEM analyses of chromium nitride in the 1.5Cu alloy aged at 1123K for 10min: (a) the bright field images, (b) Selected AreaElectron Diffraction (SAED) pattern and (c) the line analysis of Cr adjacent to chromium nitride.

0

2

4

6

8

10

Aging Time, t / min

Are

a fr

acti

on

(%

)

Base 1.5Cu

1.5Cu

Base

Chi (χ) phase

0 10 20 30

(a)

0

10

20

30

40

50

Aging Time, t / min

Are

a fr

acti

on

(%

)

Base 1.5Cu

1.5Cu

Base

Sigma (σ) phase

0 10 20 30

(b)

0 10 20 300

10

20

30

40

50

Aging Time, t / minA

rea

frac

tio

n (

%)

Base 1.5Cu

Secondary(χ+σ) phases

(c)

1.5Cu

Base

Fig. 6 Area fraction of secondary phases in the alloys aged at 1123K for 10 and 30min: (a) chi phase, (b) sigma phase and (c) secondaryphases (chi + sigma phases).

Effect of Cu on the Precipitation of Deleterious Phases and the Mechanical Properties of 27Cr­7Ni Hyper Duplex Stainless Steels 975

· phases and chromium nitride particles, were induced todecrease the total elongation of specimens. The deteriorationof ductility by the precipitation of deleterious phases wassignificantly retarded in 1.5Cu alloy compare with that inbase alloy owing to the delayed precipitation of thesedeleterious phases.

Figure 8 shows the fracture surfaces of the solution-annealed and aged base alloy and 1.5Cu alloy specimens afterthey had been subjected to tensile tests at room temperaturein air. The solution-annealed base and 1.5Cu alloy specimensexhibited fully ductile fracture surfaces with a well-developed dimple structure (Figs. 8(a) and 8(b)). However,the base alloy specimen aged at 1123K for 10min showed

somewhat complicated fractured surfaces as shown inFig. 8(c). The fracture surface showed various brittlefractured modes such as corrugated morphology,34,35) as wellas cleavage fracture36,37) and hair-line cracks.35,37) Chunet al.37) reported that cleavage fracture and hair-line crackswere deteriorated the tensile properties in high Mn austeniticsteel. Meanwhile, the fracture surface of the 1.5Cu alloyspecimen aged at 1123K for 10min showed both ductile andbrittle mixed mode (Fig. 8(d)). Although a few cleavage andintergranular fractures38) were observed in the aged 1.5 Cualloy, the area fraction of the brittle fracture surface wasgreatly reduced. Therefore, the reason that the decrease inelongation of the base alloy specimen aged at 1123K for10min (Fig. 7) was due to the easy transition of the fracturemode from ductile to brittle. Thus, this is further proof that,when added to the base alloy, Cu retards the decrease in theductility of the alloy owing to the delayed precipitation ofdeleterious phases.

4. Conclusions

(1) The addition of Cu to the base alloy reduces the totalamount of deleterious phases precipitated in it. Inparticular, Cu addition to the base alloy results inpronouncedly suppressing the amount of sigma phasewhereas it slightly facilitates the precipitation ofchromium nitride and » phase along phase boundariesand within ¡ grains.

(2) Tensile tests were used to evaluate the mechanicalproperties of the experimental alloys. The results of thetest suggested that the embrittlement of the base alloy

0 5 10 15 20 25 30 350

200

400

600

800

1000

En

gin

eeri

ng

Str

eng

th, σ

/ MP

a

Engineering Strain, ε / %

Base alloy – S.A.

1.5Cu alloy – S.A.

1.5Cu alloy – 10 min.

Base alloy – 10 min.

Fig. 7 Tensile flow curve of the solution annealed alloys and alloys aged at1123K for 10min.

10 μm

10 μm

10 μm

10 μm

2 1

2

2

3

1 : Corrugated morphology2 : Cleavage fracture3 : Hair line crack

21

1

1 : Intergranlular fracture2 : Cleavage fracture

Base alloy-solution annealed 1.5 Cu alloy-solution annealed

3

(d)(c)

Base alloy-10 min. 1.5 Cu alloy-10 min.

(a) (b)

Fig. 8 SEM image of fracture surfaces after tensile tests of the alloys: (a) solution-annealed base alloy, (b) solution-annealed 1.5Cu alloy,(c) base alloy aged at 1123K for 10min and (d) 1.5Cu alloy aged at 1123K for 10min.

S.-H. Jeon et al.976

might also be due to the precipitation of deleteriousphases such as the · and » phases and chromium nitrideparticles. The addition of Cu to the base alloy reducesits embrittlement owing to the delayed precipitation ofthese deleterious phases.

(3) The preferential precipitation of chromium nitride andthe » phase along the phase boundaries and withinthe ¡ grains seems to be closely associated with theretardation of the precipitation of the · phase. Since theformation of the · phase in the experimental alloysrequired high concentrations of Cr, Mo and W, thepreferential precipitation of chromium nitride and the »phase in the alloys in the initial stage of aging inhibitedthe nucleation and growth of the · phase. This wasbecause of the depletion of the Cr adjacent to thechromium nitride particles and the depletion of the Moand W adjacent to the » phase.

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Effect of Cu on the Precipitation of Deleterious Phases and the Mechanical Properties of 27Cr­7Ni Hyper Duplex Stainless Steels 977