transient liquid phase bonding of nitrogen containing duplex … · 2003-09-24 · transient liquid...

Transient Liquid Phase Bonding of Nitrogen Containing Duplex

Stainless Steel UNS S31803 Using Ni–Cr–Fe–Si–B Insert Metal

Byongho Rhee1;*, Sungjoo Roh2 and Dohyang Kim1

1Department of Metallurgical Engineering, Yonsei University, Seoul 120-749 Korea2Vitzrotech Co., Ltd. 605-2 Sunggok Dong, Ansan City, Gyunggi-Do, Korea

Microstructural evolution during transient liquid phase (TLP) bonding of nitrogen containing duplex stainless steel UNS S31803 has beeninvestigated. In order to evaluate mechanical property of joint, tensile strength test was carried out at room temperature. TLP bonding wasconducted at the temperature range 1283–1353K for 0–1000 s under a vacuum of 6.7mPa using Ni–7mass%Cr–3mass%Fe–4.5mass%Si–3.2mass%B amorphous insert metal. The results show that the volume fraction of austenite (�) decreased with increasing bonding temperatureand holding time. Particularly, in the case of prolonged holding time, the depleted area of � phase was observed in the base metal adjacent tojoints. There were linear correlations between the width of the remaining liquid phase and square root of holding time at each bondingtemperature. In this investigation, the secondary phases formed in the joint area were (Cr, Mo) borides dominantly. For the specimen bonded forlonger time up to 1000 s, boron nitride formed at the center and interface of joint area, on the other hand, the amount of borides decreasedcompared with the case of shorter bonding time. Tensile strength increased with holding time, and the bonding efficiency was �94% for thespecimen held for 1000 s at 1353K. Tensile strength of joint depended on, for a short holding time, brittle eutectic and borides, and aftercompletion of isothermal solidification, depended on the boron nitride formed at the joint interface.

(Received January 20, 2003; Accepted March 25, 2003)

Keywords: Transient liquid phase bonding; Duplex stainless steel; Ni–7mass%Cr–3mass%Fe–4.5mass%Si–3.2mass%B amorphous Insert

metal

1. Introduction

Although commercial austenitic stainless steels havesuperior general corrosion resistance and low temperatureimpact toughness, crevice corrosion and stress resistances arepoor in chloride-containing environments. Therefore, theapplication of the austenitic stainless steels to environmentsthat are not severely corrosive has been limited. On the otherhand, the duplex stainless steels containing both ferrite (�)and austenite (�) phases have better localized corrosionresistance than the single-phase austenitic stainless steels inchloride containing solution. In addition, as compared withthe austenitic stainless steels, the duplex stainless steels havelower cost performance because of the lower nickel content.Therefore, the application area of the duplex stainless steelsas structural materials in various industrial sectors has beenincreased steadily.1–3)

However, when inappropriate thermal cycle is adoptedduring the fabrication process, such as fusion welding, �-�ratio may be changed together with the precipitation ofintermetallic compounds (e.g. � phase, � phase, Cr23C6,Cr2N, etc.), reducing mechanical property and corrosionresistance of original materials. It is, therefore, stronglyrequired that the thermal cycle should be controlled preciselyto optimize the mechanical property and corrosion resistanceof the joints.

TLP bonding has been known as one of the methodscapable of controlling thermal cycle precisely. The processinvolves using an interlayer that melts at the bondingtemperature. With a sufficient holding time at the bondingtemperature, the interdiffusion of solutes (melting pointdepressant) at the liquid filler-base metal interface results in

isothermal solidification. Subsequently, given a suitablehomogenization treatment, the microstructure and mechan-ical properties of the joint can resemble those of the basemetal.4) Most of works on TLP bonding have beenconcentrated on Ni base superalloys.5–7) But, very fewresearch results have been reported on the TLP bonding ofduplex stainless steels, although their applications have beenincreasing. Kang et al. reported on brazing duplex stainlesssteel UNS S32550 using Ni base insert metal. They suggestedthat at the initial stage of bonding eutectic phases were notformed, and isothermal solidification process is controlled bythe formation of boron nitrides.8)

In this work, during TLP bonding of duplex stainless steelUNS S31803 using Ni–Cr–Fe–Si–B insert metal, change ofaustenite fraction in the base metal, mechanism of isothermalsolidification and formation of secondary phases wereinvestigated. Effects of microstructures of joints on themechanical property were discussed.

2. Experimental Procedure

Duplex stainless steel used in this work was UNS S31803grade in the form of bar, manufactured by AVESTA. Thesupplied duplex stainless steel bar was annealed at 1343K for10min, followed by water quenching. Chen et al. reportedthat the temperature range of 1293–1353K is suitable for thesolution treatment in a 2205 duplex stainless steel (UNSS31803) to dissolve the secondary precipitates withoutaffecting the balance of matrix phases.9) Therefore, Ni–Cr–Fe–Si–B amorphous insert metal (thickness 40 mm), appro-priate for bonding in the above temperature range and withgood corrosion resistance, was used. The compositions of thematerials in the present study are shown in Table 1.

The specimens with dimension of 10mm diameter and10mm length were machined from the bar. Their faying

*Corresponding author: Vitzrotech Co., Ltd. 605-2 Sunggok-Dong, Ansan-

City, Gyunggi-Do, Korea. E-mail: [email protected]

Materials Transactions, Vol. 44, No. 5 (2003) pp. 1014 to 1023#2003 The Japan Institute of MetalsEXPRESS REGULAR ARTICLE

surface was perpendicular to the drawing direction. Theywere ground with 1200 grit silicon carbide emery paper tomaintain the flatness of �2 mm. All specimens weredegreased with acetone in an ultrasonic cleaner aftergrinding. 40 mm spacer was inserted between the fayingsurfaces to maintain parallel and unrestrained liquid widthduring bonding (Fig. 1). TLP bonding was conducted at thetemperatures ranging of 1283–1353K for 0–1000 s under avacuum of 6.7mPa. A heating rate of 10K/s and an averagecooling rate of 20K/s from bonding temperature to 1073Kwere used.

Cross-sections of joints were prepared for microstructureexamination. The cross-sections were electrolytically etchedat 7V in 10N of NaOH solution to observe � and � phases ofthe base metal and at 4.5V in 5 vol% of H2SO4 solution toobserve the microstructure of the joints. Microstructureswere investigated using optical microscope (OM) andscanning electron microscope (SEM). Compositions wereanalyzed using energy dispersive X-ray spectroscopy (EDS)and electron probe X-ray microanalyzer (EPMA). Volumefraction of the � phase and width of the remained liquid phasewere measured using an image analyzer.

Tensile strength tests were carried out at room temperaturewith the crosshead speed of 1mm/min of speed. Fracturedsurfaces and cross-sectional microstructures were investiga-ted using OM and SEM.

3. Results and Discussion

3.1 Microstructures of base metalsFigure 2 shows the optical microstructure of the duplex

stainless steel used in the present study. The light and darkphases are � and �, respectively. The volume fraction of the �phase measured by the image analyzer was about 50%.

Figure 3 shows the optical microstructures of the basemetals bonded at 1283K and 1353K, for respectively 0 s and1000 s. Figures 4(a) and (b) show the variations of the volumefraction of � phase in the base metal adjacent to the joint

interface after bonding at 1283K and 1353K, respectively. Incase of a holding time of 0 s at 1283K, the microstructure ofbase metal is similar to original material (Fig. 3(a)). (Zerosecond means that, immediately after the bonding tempera-ture reaches the set value, heating is interrupted.) However,volume fraction of � phase decreased with increasingbonding temperature and holding time. This is because thatduring bonding a part of the original � phase was dissolved,and cooling rate was too fast to reform � phase duringcooling. Liou et al. reported that the amount of reformed �phase increased with increasing cooling time between 1073Kand 773K and the reformation of � phase is controlled bydiffusion of � stabilizing elements, such as N and Ni.10) Incase of prolonged holding time, � depleted area was observedin the base metal adjacent to joints (Figs. 3(b) and (d)). Thedepletion of the � phase can occur since � stabilizing element(N) contained in the base metal diffuses into the insert metal,while � stabilizing element (Si) diffuses into the base metalduring bonding.

3.2 Microstructures of joints and isothermalsolidification

Figure 5 shows the optical microstructures of joints at eachtemperature with the holding time of 0 s. As can be seen inFig. 5, dissolution of base metal occurred uniformly at allbonding temperatures. The rate of dissolution increased withincreasing the bonding temperature.

Figure 6 shows SEM microstructure of the remainingliquid phase after bonding at 1313K for 30 s. As can be seenin Fig. 6(b), the remaining liquid phase consisted of massivephase (marked B) and lamellar structure (marked C). Sincethe composition of insert metal used in the present study isclose to eutectic composition, the lamellar structure in Fig.6(b) is formed by eutectic reaction of �-Ni solution andintermetallic compound. The eutectic structure is known tobe detrimental to mechanical properties due to its brittle-ness.11)

The effects of holding time on the microstructures of jointsat 1313K are shown in Fig. 7. The width of the remainingliquid phase decreased with increasing holding time. Afterbonding for 1000 s, remaining liquid phase was not observed.As can be seen in Fig. 7, in spite of the holding time up to1000 s the width of the joint was nearly constant, which

Table 1 Chemical composition of materials used (mass%).

Materials B N Mn Si Ni Cr Mo Fe

Base metal UNS S31803 0.16 1.46 0.56 5.04 22.75 3.19 bal.

Insert metal MBF-20 3.2 4.5 bal. 7 3

Fig. 1 Schematic illustration of TLP bonding specimen.

Fig. 2 Optical microstructure of the base metal; duplex stainless steel

(UNS S31803).

Transient Liquid Phase Bonding of Nitrogen Containing Duplex Stainless Steel UNS S31803 Using Ni–Cr–Fe–Si–B Insert Metal 1015

0 200 400 600 800 100015

20

25

30

35

40

45

50

Distance from Base Metal/Bonded Interlayer Boundary, d/µm

Aus

teni

te P

hase

Vol

ume

Frac

tion(

%)

Aus

teni

te P

hase

Vol

ume

Frac

tion(

%)

1283K(0 s)1283K(100 s)1283K(1000 s)

0 200 400 600 800 100015

20

25

30

35

40

45

1353K(0 s)1353K(100 s)1353K(1000 s)

Distance from Base Metal/Bonded Interlayer Boundary, d/µm

50

(a) (b)

Fig. 4 Variation of austenite phase volume fraction in base metal after bonding.

Fig. 3 Optical microstructures of the base metal in the vicinity of joints; (a) 1283K–0 s, (b) 1283K–1000 s, (c) 1353K–0 s,

(d) 1353K–1000 s.

Fig. 5 Optical microstructures of the joint area; (a) 1283K–0 s, (b) 1313K–0 s, (c) 1353K–0 s.

1016 B. Rhee, S. Roh and D. Kim

indicates that the base metal dissolution is completed at theearly stage (0 s) during bonding. The same type of behaviorwas observed for the joints bonded at 1283K and 1353K.

The variations of the width of the remaining liquid phasewith holding time at each bonding temperature are shown inFig. 8. During TLP bonding, the disappearance of liquid injoints starts after the dissolution of the base metal iscompleted.12) Since the dissolution of the base metal iscompleted at the early stage (Fig. 7), the disappearance of theliquid was considered to start just after the bondingtemperature was reached. As can be seen in Fig. 8, therewere linear correlations between the width of the remainingliquid phase and square root of the holding time at eachbonding temperature. The decreasing rate of the liquid phasebecame higher with increasing bonding temperature. Itindicates that the process of isothermal solidification iscontrolled by diffusion. Generally, in TLP bonding process,isothermal solidification is controlled by diffusion of meltingpoint depressant contained in the insert metal. It can be easilyverified that isothermal solidification is controlled by diffu-

Fig. 6 SEM microstructures of the joint area after bonding at 1313K for 30 s; (a) joint area, (b) enlarged of A area marked in (a).

Fig. 7 SEM microstructures of the joint area after bonding at 1313K for; (a) 0 s, (b) 30 s, (c) 100 s, (d) 250 s, (e) 500 s, (f) 1000 s.

0 5 10 15 20 25 30 35 400

5

10

15

20

25

30

35

40

45

50

Holding Time, t/s1/2

Rem

aini

ng L

iqui

d P

hase

Wid

th, w

/µm

1283K1313K1353K

Fig. 8 Relation between the square of the holding time and the width of the

remaining liquid during isothermal solidification process.


sion of the solute element in the binary system. In multi-component system, however, the situation is more complex.Since the materials used in this work are multi-componentsystem, the isothermal solidification is controlled by complexdiffusion reaction of several elements and compound forma-tion reaction.

Meanwhile, in case that initial substance having C0

concentration exists in finite area namely, �h < x < h, theconcentration distribution with time can be expressed byeq. (1).13)

Cðx; tÞ ¼ C0=2½erfððh� xÞ=ð2ðDtÞ1=2ÞÞþ erfððhþ xÞ=ð2ðDtÞ1=2ÞÞ�

ð1Þ

where, D; diffusion coefficient, t; time, and x; distance.Here, it is assumed that the base metal is pure Fe. When

melting point depressant elements (Si, B) in insert metaldiffuse into the base metal, concentration (Cp) variation atthe center of the insert metal calculated by eq. (1) is shown inFig. 9. Here, insert metal thickness (2 h) is 40 mm, and theinitial concentrations of Si, B are 4.5mass% and 3.2mass%respectively. Diffusion coefficients of Si and B in pure Fewere calculated using vibration factor (D0Si ¼7:4� 10�5 m2/s, D0B ¼ 2:6� 10�7 m2/s) and activationenergy (QSi ¼ 220 kJ/mol, QB ¼ 95 kJ/mol).14)

From the result of EDS analysis, the average concentrationof Si in the isothermally solidified joint area adjacent to solid-liquid interface was �3:2mass%. In the case of boron,quantitative analysis was difficult due to presence of smallamount. Therefore, content of boron in �-Ni was assumed tobe 0.3 at%, which is the maximum solubility of boron in theNi–B binary system. Concentration variations of Si and B inthe joint area are shown in Fig. 9. Isothermal solidificationcompletion times calculated by the diffusion equation, andextrapolated by experimental result are compared in Table 2.Experimentally extrapolated time was shorter than thosecalculated by diffusion equation at all bonding temperature.This can be regarded because of the difference betweenactual complex reaction in multi-component system andsimple model calculation. However, in the case of calculationof Si diffusion model, isothermal solidification completion

time can be predicted more accurately by calculation than thecase of B diffusion model.

3.3 Formation of secondary phase at joint and joint-base metal interface.

Figure 10 shows SEM microstructures and element dis-tribution of the joint area bonded at 1313K for 250 s. As canbe seen in the enlarged image of A and B, The secondaryphases formed in the joint area (marked A) and in the basemetal adjacent to the joint area (marked B). EPMA analysisshowed that the lump-shaped phase in the joint area (markedA) and the rod-shaped phase in the base metal were enrichedin Cr, Mo, B elements. Therefore, it is considered that thesecondary phases are mainly (Cr, Mo) borides. The rod-shaped phase in the base metal grew along the �-� grainboundary dominantly.

On the other hand, Fig. 11 shows SEM microstructure ofthe joint area bonded at 1313K for 1000 s. Althoughisothermal solidification was completed, there were stillsecondary phases at the center of the joint. Particularly, thesecondary phase other than what were observed in the jointarea bonded at 1313K for 250 s appeared clearly in the jointarea. EPMA analysis showed that the secondary phases wereenriched in B, N elements. Therefore, it is considered that thesecondary phase is B nitride. In the vicinity of the jointinterface, there were two shape B nitrides. One close to theinsert metal is a small spherical shape, and another close tothe base metal is a massive shape.

The (Cr, Mo) boride (dark phase in Fig. 7) was continu-ously coarsened with increasing holding time up to 100 s.However, the fraction of the (Cr, Mo) boride decreased afterbonding for 100 s. Then, the B nitride at the joint interfacecoarsened with increasing holding time up to 1000 s. At thecenter of the joint, the B nitride coarsened with the shrinkageof the (Cr, Mo) boride (Fig. 11). Still, the amounts of (Cr,Mo) borides were dominant.

The mechanism of the formation of the secondary phasesin the joint area during bonding is as follows. At the earlystage of bonding, base metal is dissolved with the reaction ofbase metal and insert metal at the joint interface. Theelements contained in the dissolved base metal, such as Fe,Cr, Mo, N, etc diffuse into liquid insert metal, and theelements contained in the liquid insert metal, such as Ni, Si,B, etc diffuse into base metal. In the process, the stable phase,which is the lowest Gibbs free energy, is formed preferen-tially. Figure 12 shows the standard Gibbs energies forformation of borides and nitrides calculated by thermody-namic data.15) As can be seen in Fig. 12(a), B nitride is morestable than Cr borides. Therefore, the B nitride is formedpreferentially at the join interface. Although B nitride isstable thermodynamically, since the content of nitrogencontained in the base metal is meager, its quantity is very

0 2000 4000 6000 8000 100000

1

2

3

4

5

1353K

1313K

1283K

Boron

Cp(

mas

s%)

Holding Time, t/s

Silicon

Fig. 9 Concentration (Cp)-time curves at the center of the joint area

(calculated by eq. (1)).

Table 2 Completion time for isothermal solidification.

Bonding temp. Experimental Calculated (Si) Calculated (B)

1283K 1.53 ks 2.17 ks 12.05 ks

1313K 0.88 ks 1.36 ks 9.83 ks

1353K 0.61 ks 0.75 ks 7.60 ks


small.In the case that a specimen (the sink of solute) of

composition C1 between specimens (the source of solute) ofcomposition C0 exists in finite area namely, �h < x < h, theconcentration profile is given by eq. (2), which is convertedfrom eq. (1).

Cðx; tÞ ¼ C0 � C0=2½erfððh� xÞ=ð2ðDtÞ1=2ÞÞþ erfððhþ xÞ=ð2ðDtÞ1=2ÞÞ� þ C1

ð2Þ

where, D; diffusion coefficient, t; time, and x; distance.Here, it is assumed that the insert metal is pure Fe (C1 ¼ 0)

and there is no the dissolution of the base metal. Whennitrogen contained in base metal diffuses into the insert metalat 1313K, the concentration profile of nitrogen calculated byeq. (2) is shown in Fig. 13. The concentration profile of boronin Fig. 13 was calculated by eq. (1). Here, insert metal

thickness (2 h) is 40 mm, and the initial concentrations of N, Bare 0.63 at% and 14.50 at% respectively. Diffusion coeffi-cient of nitrogen in pure �-Fe was calculated using vibrationfactor (D0N ¼ 3:6� 10�5 m2/s) and activation energy(QN ¼ 157 kJ/mol).14)

The secondary phases forming in the joint area duringbonding can be predicted as following. As can be seen in Fig.13(a), at the early stage of bonding, nitrogen and boroncoexist only in the vicinity of the joint interface. The contentof nitrogen diffused into the insert metal is very smallcompare with boron. Therefore, the excess boron afterforming B nitride diffuses into base metal along the �-� grainboundary preferentially, and Cr borides with the reaction ofthe excess boron and Cr contained in the base metal wereformed. As can be seen in Fig. 12(b), Cr borides are morestable than Cr nitride. On the other hand, most of the phases

Fig. 10 SEM microstructures of the joint area and element profiles analyzed by EPMA.


Fig. 11 SEM microstructures of the joint area and element profiles analyzed by EPMA.

400 600 800 1000 1200 1400-50

-40

-30

-20

-10

0

BN

1/2CrB2

Temperature (K)

CrB

400 600 800 1000 1200 1400-50

-40

-30

-20

-10

0

CrB2

CrB

1/2Cr2N

Temperature (K)

(a) (b)

∆ G

f° [

kcal

/mol

, B]

∆ G

f° [

kcal

/mol

, Cr]

Fig. 12 Gibbs free energies for formation of borides and nitrides.


formed inside the joint are (Cr, Ni, etc.) borides. Here, Modetected by EPMA in Fig. 10 is not considered. Withincreasing holding time, the quantity of B nitride in the jointarea is increased. However, for a short holding time (Fig.13(b)), since the content of boron is still richer in the jointarea, the fraction of boride is dominant. For 1000 s (Fig.13(c)), the fraction of boride decreases with increasing thefraction of nitride. This coincides well with experimentalresult (Fig. 11).

Meantime, as can be seen in Fig. 11, the B nitride wasricher in the joint interface than inside joint. It is consideredthat most of nitrogen reacts with boron in the joint interfaceprior to the diffusion of nitrogen into the joint, and theisothermal solidified layer roles as the diffusion barrier ofnitrogen.

3.4 Tensile strength of jointsThe result of the tensile strength test for specimens, held

for 0 s, 100 s and 1000 s at each bonding temperature, isshown in Fig. 14. The bonding efficiency was �94% for thespecimen held for 1000 s at 1353K. Tensile strengthincreased with holding time at all bonding temperature,

since the amount of brittle eutectic and borides decreasedwith holding time. Figure 15 shows SEM microstructure ofthe joint area bonded at 1313K for 50 s. It can be seen that thecrack occurred at the center region of the joint layer. This isconsidered that the crack occurs because of stress induced byshrinkage during solidification of the remaining liquid phase.However, the crack was not observed in the joint area bondedfor more than 100 s, because of the decrement of the amountof brittle eutectic and borides. Therefore, the tensile strengthof the joint abruptly increased by prolonging the holding timemore than 100 s.

SEM microstructures of the cross section and fractographsof fractured surface after the tensile test of the specimenbonded at 1313K are shown in Fig. 16. In the case of 0 sholding time, fracture occurred at the center of the jointwhere liquid phase remained. Fracture mode was brittle-flatmode without plastic deformation. In the case of 100 sholding time, ductile dimple mode at the joint interface andbrittle-flat mode at the center of the joint was observed. In thecase of 1000 s holding time, full ductile dimple mode wasobserved at the joint interface. In this case, from SEMmicrograph and EPMA analysis, it can be seen that thefracture of the joint occurs in the place of B nitride formed atthe joint interface. These results indicate that tensile strengthof joint depended on, for a short holding time, brittle eutecticand borides, and after completion of isothermal solidification,depended on the boron nitride formed at the joint interface.

-50 -40 -30 -20 -10 0 10 20 30 40 500

2

4

6

8

10

12

14

16B MB M I M

Nitrogen

Con

cent

ratio

n, a

t%

Distance, d/µm

Boron

-300 -200 -100 0 100 200 3000.0

0.5

1.0

1.5

2.0

2.5

Distance, d/µm

Con

cent

ratio

n, a

t%

Nitrogen

Boron

-300 -200 -100 0 100 200 3000.0

0.5

1.0

1.5

2.0

2.5

Distance, d/µm

Con

cent

ratio

n, a

t%

Nitrogen

Boron

(a) (b) (c)

Fig. 13 Nitrogen and boron distribution calculated by eqs. (1) and (2) in BM(Base Metal)-IM(Insert Metal)-BM couple at 1313K for;

(a) 0.1 s, (b) 100 s, (c) 1000 s.

0 200 400 600 800 10000

100

200

300

400

500

600

700

800

900

1000

Base Metal

Tens

ile S

tren

gth,

σ/M

Pa

Holding Time, t/s

1283K1313K1353K

Fig. 14 Tensile strength test results of the bonded materials for the

different temperature and holding time.

Fig. 15 SEM microstructure of the joint area bonded at 1313K for 50 s.


Fig. 16 SEM microstructures and element profile of fractured surface after tensile strength test.


4. Conclusions

(1) The volume fraction of � phase in the base metaldecrease with increasing bonding temperature andholding time. Particularly, in the case of prolongedholding time, the depleted area of � phase is observed inthe base metal adjacent to joints. It is consideredbecause � stabilizing element (N) contained in the basemetal diffuses into the insert metal, while � stabilizingelement (Si) diffuses into the base metal duringbonding.

(2) Isothermal solidification has a linear correlation withholding time1=2. The isothermal solidification behavioroccurs with complex reactions of constituents, but ismainly controlled by diffusion of melting point depres-sant elements. Especially, in the case of calculation ofSi diffusion model, isothermal solidification completiontime can be predicted more accurately by calculationthan the case of B diffusion model.

(3) In this investigation, the secondary phases formed in thejoint area are (Cr, Mo) borides dominantly. For thespecimen bonded for longer time up to 1000 s, boronnitride forms at the center and interface of joint area, onthe other hand, the amount of borides decreasedcompared with the case of shorter bonding time. It isconsidered that the formation of boron nitride dependson diffusion of nitrogen and boron. Borides formed injoint-base metal interface grow along �-� grain bound-ary preferentially.

(4) Tensile strength increases with holding time up to1000 s at all bonding temperature. The highest bondingefficiency is �94% (for the specimen held for 1000 s at1353K). Tensile strength of joint depends on, for ashort holding time, brittle eutectic and borides, and aftercompletion of isothermal solidification, depends on theboron nitride formed at the joint interface.

REFERENCES

1) J. O. Nilsson: Mater. Sci. Technol. 8 (1992) 685–700.

2) H. D. Solomon and T. M. Devine: R. A. Lura (Ed), Duplex Stainless

Steel, ASM (1982) pp. 693–756.

3) N. Lopez, M. Cid and M. Puiggali: Corros. Sci. 41 (1999) 1615–1631.

4) T. I. Khan and E. R. Wallach: J. Mater. Sci. 31 (1996) 2937–2943.

5) H. Ikawa and Y. Nakao: Trans. Japan Welding Society 8 (1977) 3–8.

6) Y. Nakao, K. Nishimoto, K. Shinozaki and C. Y. Kang: Proceedings

Superalloys ’89, (Metall. Soc. AIME, 1989) 775–783.

7) W. F. Gale and S. V. Orel: Metall. Trans. A 27A (1996) 1925–1931.

8) C. Y. Kang, C. S. Woo, T. K. Han, C. H. Hong, D. U. Kim and Y. C.

Woo: Proc. Conf. ‘2001 International Brazing & Soldering’, Yangz-

hong, China, 148–155.

9) T. H. Chen and J. R. Yang: Mater. Sci. Eng. A 311 (2001) 28–41.

10) H. Y. Liou, R. I. Hsieh and W. T. Tsai: Corrosion Science 44 (2002)

2841–2856.

11) X. W. Wu, R. S. Chandel, H. P. Seow and H. Li: J. Mater. Proc.

Technol. 113 (2001) 215–221.

12) K. Saida, Y. Zhou and T. H. North: J. Japan Inst. Metals 58 (1994) 810–

818.

13) J. Crank: The Mathematics of Diffusion, (Oxford University Press,

1975) pp. 13–16.

14) The Japan Institute of Metals: Metals Data Book, (Maruzen, 1993)

pp. 21–22.

15) O. Kubaschewski and C. B. Alcock: Metallurgical Thermochemistry,

5th ed, (Pergamon Press, 1979) pp. 336–356.


transient liquid phase bonding of nitrogen containing duplex … · 2003-09-24 · transient liquid...

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