Mechanism of Intermediate Temperature Embrittlement of Ni and Ni-based Superalloys

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Mechanism of Intermediate TemperatureEmbrittlement of Ni and Ni-based SuperalloysLei Zheng a b , Guido Schmitz b , Ye Meng a , Reda Chellali b & Ralf Schlesiger ba School of Materials Science and Engineering , University of Science and TechnologyBeijing , Beijing , P.R. Chinab Institute of Materials Physics , University of Muenster , Muenster , GermanyPublished online: 10 Sep 2012.

To cite this article: Lei Zheng , Guido Schmitz , Ye Meng , Reda Chellali & Ralf Schlesiger (2012) Mechanism of IntermediateTemperature Embrittlement of Ni and Ni-based Superalloys, Critical Reviews in Solid State and Materials Sciences, 37:3,181-214, DOI: 10.1080/10408436.2011.613492

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Critical Reviews in Solid State and Materials Sciences, 37:181–214, 2012Copyright c© Taylor and Francis Group, LLCISSN: 1040-8436 print / 1547-6561 onlineDOI: 10.1080/10408436.2011.613492

Mechanism of Intermediate Temperature Embrittlementof Ni and Ni-based Superalloys

Lei Zheng,1,2,∗ Guido Schmitz,2 Ye Meng,1 Reda Chellali,2

and Ralf Schlesiger2

1School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing,P.R. China2Institute of Materials Physics, University of Muenster, Muenster, Germany

Ni-based superalloys play an important role in aircraft engine propulsion. However, manyexperiments demonstrated that these alloys as well as Ni always show ductility loss at interme-diate temperature, which constrains further development. A comparison of published papersby various authors reveals considerable differences in understanding the mechanism of inter-mediate temperature embrittlement. To clarify this situation, the present article first confirmsthe generality of intermediate temperature embrittlement of Ni and Ni-based alloys by theexperimental results reported in the literature. The existing interpretations of the mechanismare then outlined. Based on the generality, these interpretations are discussed through therepresentative investigations on intermediate temperature embrittlement. It is shown that themechanism of intermediate temperature embrittlement has not been satisfactorily explainedyet and “nonequilibrium interface segregation” of impurities taking into account the effect ofstrain rate may be the origin of intermediate temperature embrittlement of Ni and Ni-basedsuperalloys. Future research directions aiming at the reason of abnormal fracture mode, theeffect of the state of applied stress, the influence of strain rate, and the development of the theorynonequilibrium grain boundary segregation, are suggested to provide a complete understandingof intermediate temperature embrittlement.

Keywords Ni-based superalloy, intermediate temperature embrittlement, grain boundary segrega-tion, fracture mode, applied stress, strain rate

Table of Contents

1. INTRODUCTION .............................................................................................................................................. 182

2. GENERALITY OF INTERMEDIATE TEMPERATURE EMBRITTLEMENT OF Ni AND Ni-BASED SUPER-ALLOYS ............................................................................................................................................................. 1832.1. Intermediate Temperature Embrittlement of Wrought Ni-Based Superalloys ...................................................... 1832.2. Intermediate Temperature Embrittlement of Cast Ni-Based Superalloys ............................................................ 1842.3. Intermediate Temperature Embrittlement of Powder Metallurgy Ni-Based Superalloys ....................................... 1842.4. Intermediate Temperature Embrittlement of Highly and Commercially Pure Ni ................................................. 1852.5. Summary ..................................................................................................................................................... 185

3. THE EXISTING INTERPRETATIONS FOR INTERMEDIATE TEMPERATURE EMBRITTLEMENT .......... 1863.1. The Interpretation due to Intergranular Precipitates .......................................................................................... 1873.2. The Interpretation due to Grain Boundary Shearing or Sliding .......................................................................... 1873.3. The Interpretation due to Gas Phase Embrittlement ......................................................................................... 1873.4. The Interpretation due to Glide Plane Decohesion ........................................................................................... 189

∗E-mail: zhenglei [email protected]

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182 L. ZHENG ET AL.

3.5. The Interpretation due to Dynamic Strain Aging ............................................................................................. 1903.6. The Interpretation due to Grain Boundary Segregation ..................................................................................... 191

3.6.1. Temperature Dependence of Intermediate Temperature Embrittlement by Equilibrium Gain BoundarySegregation ........................................................................................................................................ 191

3.6.2. Temperature Dependence of Intermediate Temperature Embrittlement by Nonequilibrium Grain BoundarySegregation ........................................................................................................................................ 192

3.7. Summary ..................................................................................................................................................... 193

4. DISCUSSIONS ................................................................................................................................................... 1934.1. Discussion on the Interpretation due to Intergranular Precipitates ...................................................................... 1944.2. Criticism on the Interpretation due to Grain Boundary Shearing or Sliding ........................................................ 1964.3. Remark on the Interpretation due to Gas Phase Embrittlement .......................................................................... 1984.4. Evaluation of the Interpretation due to Glide Plane Decohesion ........................................................................ 1994.5. Judgment of the Interpretation due to Dynamic Strain Aging ............................................................................ 2004.6. Comment on the Interpretation due to Grain Boundary Segregation .................................................................. 202

4.6.1. Equilibrium Grain Boundary Segregation ............................................................................................ 2024.6.2. Nonequilibrium Grain Boundary Segregation ....................................................................................... 202

4.6.2.1. Arguments in favor of nonequilibrium grain boundary segregation theory ................................. 2024.6.2.2. Extended application of nonequilibrium grain boundary segregation theory .............................. 203

4.7. Summary ..................................................................................................................................................... 205

5. FUTURE DIRECTIONS .................................................................................................................................... 205

6. CONCLUSION .................................................................................................................................................. 208

ACKNOWLEDGMENT ............................................................................................................................................ 208

REFERENCES ......................................................................................................................................................... 208

1. INTRODUCTIONNi-based superalloys are advanced structural materials that

are mainly used in aircraft engines and also extensively ap-plied in chemical, petrolic and electrical industries.1–30 By virtueof steady improvements in composition and processing, theirelevated-temperature properties have been evolving for morethan half century. However, plenty of experiments30–42 con-firmed that these alloys retain an embrittling behavior in in-termediate temperature range (between 500 and 900 ◦C), whichmeans that the maximum elongation or reduction in area atintermediate temperatures are much lower than those at low(≤500 ◦C) and high temperatures (≥900 ◦C) in tensile tests.The embrittlement of Ni and its alloys (Inconel and Monel) areshown in Figure 1.31 It is evident that Ni and its alloys havea ductility minimum that occurs around 600 ◦C. Nowadays,more efficient and longer life engines require the developmentof more reliable and higher temperature Ni-based superalloys.Since superalloys are operated generally in the correspondingtemperature range,1–42 intermediate temperature embrittlementposes a serious threat on the reliability of Ni-based superalloysand becomes an urgent problem to be solved.

The intermediate temperature embrittlement of metals andalloys has been an area of intensive research since 1912.43

Experimental results have demonstrated that this embrittlingphenomenon exists not only in Ni-based alloys,30–42 but alsoin Fe-based,44 Co-based,45 Cu-based,46 and Al-based alloys.47

The embrittlement of Ni and its alloys was first investigated byMerica and Waltenberg48 in 1925. In 1975, Fiore49 pointed outthat the intermediate temperature embrittlement is common toa wide range of Ni-based superalloys and plays an importantrole in cracking tendency which makes welding and deforma-tion difficult. He attributed the embrittlement to the interactionsbetween factors leading to ductility loss and those leading toductility recovery. Since then, numerous experimental resultsand theoretical analyses had been reported. However, the exist-ing interpretations for intermediate temperature embrittlementgiven by different authors are contradictory to each other. Thissituation necessitates a clear understanding of the nature of in-termediate temperature embrittlement.

This review is therefore carried out to discuss the validityof these interpretations and to figure out the mechanism ofintermediate temperature embrittlement. First, the generalityof intermediate temperature embrittlement of Ni and Ni-basedsuperalloys is confirmed. The existing interpretations are thenoutlined and summarized. Afterwards, these interpretations arediscussed in view of the related experiments and the remarkable

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INTERMEDIATE TEMPERATURE EMBRITTLEMENT 183

FIG. 1. Maximum relative elongation vs. temperature forMonel, Inconel and Ni after data from Rines and Wray.31

generality. Finally, directions of future investigations are of-fered for an overall understanding of intermediate temperatureembrittlement.

GLOSSARY� = grain boundaryI = solute atomM = solvent atomC� = equilibrium concentration of solute atoms at grain

boundaryCB = bulk concentration of solute atom�G = molar Gibbs energy of grain boundary segregation

of solute atom�G0 = standard (ideal) molar Gibbs energy of grain bound-

ary segregation of solute atom�GE = excess (interaction) molar Gibbs energy of grain

boundary segregation of solute atomµ0,�

i = standard chemical potentials of the elements (i = Iand M) in pure state at grain boundary at the temper-ature, pressure, and structure of the solvent M

µ0i = standard chemical potentials of the elements (i = I

and M) in pure state in the bulk at the temperature,pressure, and structure of the solvent M

γ �i = activity coefficients of the elements (i = I and M) at

grain boundaryγ i = activity coefficients of the elements (i = I and M) in

the bulkT0 = a higher temperatureT = a certain lower temperatureC�(t) = grain boundary concentration of solute atom after

time t at a temperature TC�(0) = initial grain boundary concentration of solute atom

tc = critical time in grain boundary segregation of soluteatom at a temperature T

D = volume diffusion coefficient of solute atom at a tem-perature T

α = equilibrium ratio C�/CB

d = grain boundary thicknessr = average grain radiusδ = numerical factorDc = diffusion coefficient of solute atom-vacancy com-

plexes in matrixCm

�(T0) = maximum grain boundary concentration of soluteatom at temperature T0

Cm�(T) = maximum grain boundary concentration of solute

atom at a temperature TEb = binding energy between vacancy and solute atomEf = vacancy formation energyK = Boltzmann’s constantαN = ratio Cm

�(T)/CB

C�(tc) = grain boundary concentration of solute atom at thecritical time tc

σ 1 = maximum principal stressσ 2 = principal stressσ 3 = minimum principal stressµσ = Rod parameterαS = soft coefficientRσ = stress triaxiaty

2. GENERALITY OF INTERMEDIATE TEMPERATUREEMBRITTLEMENT OF Ni AND Ni-BASEDSUPERALLOYSNi-based superalloys can be classified into wrought, cast

and powder metallurgy alloys according to manufacturingroutes,1–6,28–30 and into solution-strengthened and age-hardenedalloys according to strengthening mechanism.1–6,28–30 In thissection, the intermediate temperature embrittlement of Ni-basedsuperalloys is displayed in terms of manufacturing routes. Theembrittlement of highly and commercially pure Ni is exhibitedas well.

2.1. Intermediate Temperature Embrittlement ofWrought Ni-Based Superalloys

Wrought superalloys are the most widely used among threetypes of Ni-based superalloys.1–6,28–30 They possess differentmicrostructures and retain different mechanical properties bymeans of various heat treatments and plastic deformations.

In 1961, Rhines and Wray31 first reported that Niand Ni-based superalloys reveal maximum embrittlementaround 600 ◦C. Since then, Inconel 718 (see Fig-ure 2),35 Udimet 115,36 Alloy 625,34,50,51 Inconel 740,52

Nimonic 263,52 Inconel 738LC,33,53 Alloy C-276,32 AlloyC-22,54 Waspaloy,55 Hastelloy X (see Figure 3),56–58 Inconel600,59 and Ni-Cr-W alloy41 have been reported to exhibit

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184 L. ZHENG ET AL.

FIG. 2. Reduction in area and elongation vs. temperature ofwrought Inconel 718 after data from Barker.35

intermediate temperature embrittlement as well. Among thesealloys, Alloy C-276,32 Alloy C-22,54 Hastelloy X,56–58 Inconel600,59 and Ni-Cr-W alloy41 are solution-strengthened while theremaining are age-hardened.

2.2. Intermediate Temperature Embrittlement of CastNi-Based Superalloys

Important components of aircraft engines, for example tur-bine blades, are usually made by casting process.1–4 Turbineblades have polycrystalline dendritic microstructure if they areproduced by conventional casting.60–62 To improve the perfor-mance of superalloys, directional solidification is applied toeliminate transverse grain boundaries perpendicular to the prin-cipal stress axis and to achieve columnar grain structure.63–65 Tofurther improve the properties of superalloys, single crystal wasintroduced because of the high incipient melting and gammaprime solvus temperature.66–68

Conventionally cast superalloys, such as Inconel 718 (seeFigure 4),35 CM247LC,37,38 M963,69 IN939,70 MAR-M247,40

M951,42 and B190039 were demonstrated to show intermedi-

FIG. 3. Tensile properties vs. temperature of Hastelloy X.56

Tensile samples were homogenized at 1149 ◦C for 2 hoursand water quenched. Elevated temperature tensile tests werethen performed. (Reprinted with permission from Arkoosh andFiore56 Copyright 1972: Springer Science + Business Media.)

ate temperature embrittlement evidently. Directionally solidifiedsuperalloys, such as GTD-111,71–74 PWA 664 (see Figure 5),75

PWA 659 (see Figure 5),75 DZ951,76,77 and CM247LC37,38 alsodisplay obvious intermediate temperature embrittlement. Re-markably, the single crystal superalloys, PW 1484 (see Fig-ure 6),78 DD3,79 CMSX-4,80,81 and CMSX-1082 were proven toreveal intermediate temperature embrittlement, too.

2.3. Intermediate Temperature Embrittlement of PowderMetallurgy Ni-Based Superalloys

In 1970s, efforts to increase the performance of gas turbinedisks led to the development of several alloys that contain un-precedentedly high level of γ ’ phase. However, large elemen-tal contents of γ ’ formers, such as Al, Ti, and Nb, can lead toexcessive macro- or/and micro-segregation during solidification

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FIG. 4. Reduction in area and elongation vs. temperature ofconventionally cast Inconel 718 after data from Barker.35

and consequently induce significant embrittlement.83 Powdermetallurgy combining hot isostatic pressing was therefore in-troduced to form turbine disks, since this technique, besideseliminating segregation effects, offers the potential for produc-ing a fine grained structure.83,84 Moreover, significant reductionin finished component cost can be realized due to reduced ma-chining labor and improved material utilization.

The powder metallurgy superalloys, such as Inconel 71885

(see Figure 7) and Nimonic APK1,86 were confirmed to possessintermediate temperature embrittlement.

FIG. 5. Elongation of directionally solidified PWA 664 andPWA 659 after data from Piearcey and Versnyder.75

FIG. 6. Elongation of PWA1484 Ni-based single crystal su-peralloy after data from Cetel and Duhl.78 The heat treatmentbefore tensile test was: 1316 ◦C/4 hours + 1079 ◦C/4 hours +704 ◦C/24 hours.

2.4. Intermediate Temperature Embrittlement of Highlyand Commercially Pure Ni

In 1961, Bieber and Decker87 proved that highly pure Ni ap-pears to have not intermediate temperature embrittlement, whilecommercially pure and impurities added Ni reveal obvious em-brittling behaviors (see Figure 8). In 1964, Kraai and Floreen88

comprehensively studied the effect of 1 to 50 ppm S on tensileductility of Ni at 800 ◦F (427 ◦C) to 1400 ◦F (760 ◦C). Theyfound that the ductility is excellent at very low S levels (near0 ppm) and no evidence of intermediate temperature embrittle-ment is apparent, while the specimens at S levels on the orderof 5 to 20 ppm show evident ones (see Figure 9). Further ob-servations by transmission electron microscopy did not revealany second phase particles at Ni grain boundaries.89,90 Besides,commercially pure Ni27091 and Ni20092,93 were proven to pos-sess intermediate temperature embrittlement. In 2011, tensileductility of highly pure Ni and binary Ni(Bi) alloy were studiedcomparatively in the temperature range of 400 ◦C to 850 ◦C byZheng et al.30 Ni(Bi) alloy with 25 wt ppm Bi was confirmedto attain a clear maximum embrittlement between 700 ◦C and750 ◦C, while highly pure Ni does not (see Figure 10). Ob-servations by transmission electron microscopy verified that nointergranular precipitates formed in Ni(Bi) alloy at 750 ◦C.30

2.5. SummaryFrom the preceding information with respect to Ni alloys

displaying intermediate temperature embrittlement, it is obviousthat the extent of intermediate temperature embrittlement variesfrom alloy to alloy and even with heat treatment of a given alloy,but there is no doubt that the phenomenon is generally presentin these alloys.

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FIG. 7. (a) Reduction in area (RA) and (b) elongation of as- hotisostatic pressed superalloy 718 as a function of temperature.85

Parameters of hot isostatic press was 1200 ◦C/120MPa/3 hours.Tensile tests were performed at different temperatures afterhot isostatic press. (Reprinted with permission from Appa Raoet al.85 Copyright 2006: Elsevier.)

Besides, three conclusions can be drawn. First, intermedi-ate temperature embrittlement exists not only in wrought butalso in cast as well as in powder metallurgy Ni-based super-alloys. So, intermediate temperature embrittlement is indepen-dent of the manufacturing process. Second, intermediate tem-perature embrittlement exists in both solution-strengthened andage-hardened superalloys. Thus, it is not limited to a certainstrengthening mechanism. Third, commercially pure or impu-rity added Ni without precipitates can also display a clear maxi-mum of embrittlement. This indicates that contents of impuritiesin the level of ppm are sufficient to produce an embrittlement.Therefore, intermediate temperature embrittlement doesn’t rootin the evolution of second phases in Ni and Ni-based superal-loys, although second phases can affect the depth and the width

FIG. 8. The malleability of Ni.87 A) Highly pure Ni; B) Com-mercially pure Ni; C) Ni with harmful elements added; D) Ma-terial of B or C with malleabilizer added.

of intermediate temperature embrittlement as well as the tem-perature at which the maximum embrittlement appears.

3. THE EXISTING INTERPRETATIONS FORINTERMEDIATE TEMPERATURE EMBRITTLEMENTTo our knowledge, the existing interpretations to eluci-

date intermediate temperature embrittlement are based onintergranular precipitates,34,50,51,85,94 grain boundary shearing orsliding,31,41,56,59,92,93,95,96 gas phase embrittlement,86,91,97,98

FIG. 9. Ductility of Ni as a function of temperature at variousS levels.88 The tensile samples were annealed at either 1093 ◦Cor 871 ◦C to produce different grain sizes and held for 45 minat each temperature before testing.

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FIG. 10. Elongations of highly pure Ni and Ni(Bi) alloy at dif-ferent temperatures.30 Highly pure Ni and Ni(Bi) alloy wereheat-treated for 1000 ◦C/0.5 hour and then quenched in icedwater. Tensile tests were then performed. (Reprinted with per-mission from Zheng et al. Copyright 2011: Elsevier.) (Colorfigure available online).

decohesion of glide plane,36–40,42,53,69,71,72,76,77 dynamic strainaging,32,54,55 and grain boundary segregation.30,57,58,88,99 In thissection, these interpretations are outlined.

3.1. The Interpretation due to Intergranular PrecipitatesIntergranular precipitates were the first factor considered to

cause embrittlement of Ni. In 1925, Merica and Waltenberg48 al-ready attributed the detrimental effect of S on the ductility of Nito the formation of a Ni-Ni3S2 eutectic film at grain boundaries.

At lower temperature, only a few precipitates of small sizeexist at grain boundaries (Figure 11a). These small precipitateshave little effect on the initiation of a crack and therefore thealloys retain high ductility. At intermediate temperature, in-tergranular precipitates become larger, thereby leading to easyinitiation of a crack (Figure 11b) and low ductility of alloys.At higher temperature, the ductility is restored owing to thedissolution of precipitates (Figure 11c).

3.2. The Interpretation due to Grain Boundary Shearingor Sliding

This interpretation was first offered by Rhines and Wray31

based upon the mechanism of creep in crystalline substances.At low temperature, fracture propagates by normal trans-

granular crack mode (Figure 12a), thus causing a high ductility.As grain boundary shearing becomes possible at slightly highertemperature, grain boundary triple junctions (Figure 12b1) andany irregularities, such as grain boundary directional changesand steps (Figure 12b2), or intergranular precipitates (Figure12b3) will cause strain concentrations. These centers of localstrain concentrations will serve as sites for the initiation of voids.Growth and linkage of these voids induce microcracks, therebyresulting in low ductility. Continued increase in temperature ini-tializes the recrystallization (Figure 12c), which removes grainboundaries where microcracks are formed, thus deterring theirfurther growth and re-establishing high ductility.

3.3. The Interpretation due to Gas Phase EmbrittlementA considerable amount of experiments86,91,97,98,100–102 had

demonstrated that a wide range of Ni-based superalloys aresusceptible to brittle intergranular failure with reduced ductilitywhen tensile test is carried out at intermediate temperature inaggressive environment, oxygen penetration in particular.91

At low temperature, little grain boundary sliding and oxy-gen penetration appears (Figure 13a). Thus, samples behave

FIG. 11. Schematic illustration of the interpretation due to intergranular precipitates.

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FIG. 12. Schematic illustration of the interpretation due to grain boundary shearing or sliding.

normally and exhibit high ductility. When the temperature israised, oxygen penetration along grain boundary becomes sig-nificant (Figure 13b1). In consequence, grain boundaries arepinned by oxygen penetration and boundary sliding cannot oc-cur, thus causing embrittlement of samples. In another interpre-tation, grain boundaries are weakened by adsorption of oxygen

and cracks are initiated easily by stress concentrations ahead ofdislocation pile-ups at grain boundaries and/or other obstacles(Figure 13b2), which finally results in low ductility as well.Besides, oxygen penetration along grain boundaries can inducesubsequent oxidation of grain boundary particles (Figure 13b3),possibly carbides, sulfides and delta phase. Thus, continuous

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FIG. 13. Schematic illustration of the interpretation due to gas phase embrittlement.

oxides or oxide particles pin the boundaries, making grainboundary sliding or migration impossible, eventually causinggrain boundary cracking and significant loss in ductility. At hightemperature, the boundaries gain sufficient energy to overcomethe barriers of pinning. As a result, dynamic recrystallizationoccurs and then ductility recovers (Figure 13c).

With increasing moisture level in the test atmosphere, hydro-gen will become the dominant embrittling species.97 Avoiding

pointing out a specific gas, Woodford98 defined the brittle grainboundary fracture caused by any gaseous species as gas phaseembrittlement.

3.4. The Interpretation due to Glide Plane DecohesionThis interpretation was first proposed by Jensen and Tien36

in 1985. They suggested that homogeneous or localized

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FIG. 14. Schematic illustration of the interpretation due to glide plane decohesion.

deformation is the main factor affecting the variation in duc-tility with temperature.

At low temperature, the deformation is inhomogeneous andrestricted to a few slip bands (Figure 14a). With increasing tem-perature, the mobility of dislocations increases and promotesthe inhomogeneity in localized strains since more dislocationsare trapped at the γ /γ ′ interface. Continuous movement of dis-locations enhances consequent stress concentration at the γ /γ ′

interface, finally leading to glide plane decohesion (Figure 14b).In the high temperature range, the process of thermally activateddislocation climb becomes the controlling mechanism (Figure14c) and therefore deformation tends to become homogeneous.So, the ductility minimum is obtained at intermediate tempera-tures.

3.5. The Interpretation due to Dynamic Strain AgingThe behavior of dynamic strain aging can be characterized

by serrated flow in stress-strain curves.32,51,54,55,103–108 This par-ticular mode of plastic flow is induced by interaction of dis-locations with diffusing solute,32,51,54,55,103–108 which impedesthe movement of dislocations within a susceptible temperatureregime.

At low temperature, the mobility of solute atoms is muchslower than that of dislocations. Therefore, samples deformin the normal way of multiplication and movement of dislo-cations, accordingly revealing relatively high ductility (Figure15a). At intermediate temperatures, solute atoms dynamicallycluster around migrating dislocations and form Cottrell atmo-spheres due to the accelerated mobility of solute atoms. Cottrell

FIG. 15. Schematic illustration of the interpretation due to dynamic strain aging.

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atmospheres can impede the migration of dislocations, thus re-sulting in reduced failure strain (Figure 15b). However, clus-tering tendency decreases with further increasing temperatures.When the dislocations are freed from the impediment of soluteatoms at high temperature, the magnitude of failure strain andthus high ductility are restored (Figure 15c).

3.6. The Interpretation due to Grain BoundarySegregation

It is noteworthy that fracture in the temperature range ofductility minimum is frequently reported to be intergranu-lar,32,49,57,58,91,97–99 and grain boundary segregation of impuritiesplays an important role in intergranular failure.30,99,109–111 Manyexperiments had confirmed grain boundary segregation of im-purities in Ni-based superalloys112–126 and commercially pureNi.126–132 Thus, researchers correlated the intermediate temper-ature embrittlement with the grain boundary segregation of dele-terious elements.30,57,58,99

3.6.1. Temperature Dependence of Intermediate TemperatureEmbrittlement by Equilibrium Gain BoundarySegregation

Grain boundary segregation can be classified into equilib-rium and nonequilibrium types.133 McLean134 first proposed amodel for equilibrium grain boundary segregation and derivedappropriate thermodynamic and dynamic equations. The ther-modynamic equation is

C�

1 − C�

= CB

1 − CB

exp

(−�G

RT

)[1]

where C� is the equilibrium concentration of solute atoms atgrain boundary �, CB is the bulk concentration of solute atom,and �G is the molar Gibbs energy of grain boundary segregationof solute atom. �G includes the standard (ideal) contribution,�G0, and the excess (interaction) contribution, �GE.135,136 Fora binary system with the segregation of solute atom I at grainboundary of solvent element M, �G0 and �GE are definedas135,136

�G0 =(µ

0,�I (M) + µ0

M

)−(µ

0,�M + µ0

I (M)

)[2]

where µ0,�i and µ0

i denote the standard chemical potentials ofthe elements I and M in pure state at grain boundary and in thebulk at the temperature, pressure, and structure of the solventM.

�GE = RT ln

(γ �

I γM

γIγ�M

)[3]

where γ �i and γ i represent the activity coefficients of the ele-

ments I and M at grain boundary and in the bulk, respectively.The dynamic equation of equilibrium grain boundary segrega-

FIG. 16. Dynamics of equilibrium grain-boundary segrega-tion.134

tion is

C� (t) − C� (0)

C� − C� (0)= 1 − exp

(4Dt

α2d2

)· erfc

(2√

Dt

αd

)[4]

where C�(t) is the grain boundary concentration after time t at acertain temperature, C�(0) is the initial grain boundary concen-tration, D is the volume diffusion coefficient at lower tempera-ture, α is the equilibrium ratio C�/CB, and d is the grain bound-ary thickness. Figure 16 illustrates the change in grain bound-ary concentration with time from initial concentration C�(0)to final equilibrium concentration C�. Later, Guttmann,137,138

Lejcek et al.,135,136,139–141 and Kirchheim142,143 contributed to thedevelopment of equilibrium grain boundary segregation theorysignificantly.

At low temperature, the diffusion coefficients of impuritiesare too slow to induce significant grain boundary segregation ofimpurities within a given time (Figure 17a). This produces littleeffect on the ductility. As temperature increases, the diffusivityof impurities becomes accelerated. Therefore, high concentra-tions of impurities will be gathered at grain boundaries andconsequently the alloys become brittle (Figure 17b). With a fur-ther increase in temperature, the diffusion coefficients becomeeven faster. However, grain boundary segregation of impuritiesat equilibrium state, C�, decreases simultaneously according toEq. (1). As a result, grain boundary segregation of impuritiesafter a given time, C�(t), reduces according to Eq. (4) (Figure17c1) and high ductility is restored.30 In another interpreta-tion,99 ductility recovery at high temperatures was attributed todynamic recrystallization (Figure 17c2), which can release lo-cal stress concentration and hinder the linkage of intergranularcracks.

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FIG. 17. Schematic illustration of the interpretation due to equilibrium grain boundary segregation.

3.6.2. Temperature Dependence of Intermediate TemperatureEmbrittlement by Nonequilibrium Grain BoundarySegregation

The theory of nonequilibrium grain boundary segre-gation was proposed by Aust et al.144 and Anthony.145

Based on the investigation of Williams et al.,146 a kineticmodel and related equations were given by Faulkner147 andXu et al.,148–150 which had been approved by many re-searchers.114,115,151–155 Nonequilibrium grain boundary segrega-tion includes thermo-induced,147–155 irradiation-induced,156–159

stress-induced,149,160–165 and deformation-induced types.136,166

At present, the theory of thermo-induced nonequilibrium grainboundary segregation is the most developed.

It is suggested that the mechanism of thermo-inducednonequilibrium grain boundary segregation is based on a lo-cal equilibrium in which a sufficient quantity of vacancy-soluteatom complex exists. Solute atom, vacancy and their recom-

bined complex are in equilibrium with each other.146–155 Whena specimen is quickly cooled to a lower temperature from ahigher one, equilibrium concentration of vacancies can be real-ized immediately along grain boundaries. Accordingly the con-centration of the complex decreases at grain boundaries while inregions remote from grain boundaries, oversaturated vacancieswill recombine with solute atoms, making the complex concen-tration increase. Thus, a gradient in the complex concentrationis established. This gradient drives complexes to diffuse to-wards grain boundaries, thereby causing excessive solute atomsto concentrate at grain boundaries and resulting in a signifi-cant nonequilibrium segregation level. However, when the grainboundary concentration of solute atoms becomes higher than itsequilibrium value, solute atoms will start to diffuse back intograin interiors. At the beginning, the diffusion of complexes isprevailing but decreases with time, while the reverse diffusionof solute atoms increases as time further extends. Therefore,

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a critical time exists, at which the reverse diffusion of soluteatoms balances the diffusion of complexes toward the grainboundary. At this time, concentration of solute atoms reachesa maximum.147–155 If the annealing time at a given temperatureis shorter than the critical time, the diffusion of complexes isthe dominant process and referred to as segregation process. Ifthe annealing time is longer than the critical time, the reversediffusion of solute atoms is dominant and referred to as de-segregation process. Naturally, themo-induced non-equilibriumgrain boundary segregation disappears as time approaches in-finity and distribution of atomic species reach full equilibrium.The critical time, tc, is given by147–155

tc = r2 ln (Dc/D)

4δ (Dc − D)[5]

where r is the average grain radius, δ is a numerical factor, andDc is the diffusion coefficient of complexes in matrix. When asample is cooled from a higher temperature, T0, to a lower one,T , the maximum grain boundary concentration, Cm

�(T), is givenby147–149,153,154

Cm� (T ) = CB

(Eb

Ef

)exp

[(Eb − Ef

kT0

)−(

Eb − Ef

kT

)][6]

where Eb is the binding energy between vacancy and soluteatom, Ef is the vacancy formation energy, and k is the Boltz-mann’s constant. Eq. (6) is an important expression that de-scribes the level of thermo-induced nonequilibrium grain bound-ary segregation. On the basis of this equation, kinetic equationsof segregation and desegregation were derived by Xu et al.148–151

Suppose that the sample is cooled so quickly from temperatureT0 down to T that no mass transfer occurs in the specimen dur-ing cooling, a dynamic equation of the segregation process canbe formulated as

C� (t) − Cm� (T0)

Cm� (T ) − Cm

� (T0)= 1 − exp

(4Dct

α2Nd2

)erfc

[2 (Dct)1/2

αNd

][7]

where C�(t) is the grain boundary concentration after time tat temperature T , Cm

�(T0) and Cm�(T) are the maximum grain

boundary concentrations at temperatures of T0 and T , and αN isthe ratio Cm

�(T)/CB. The dynamic equation of the desegregationprocess is formulated as

C� (t) − CB

C� (tc) − CB

= erf

{d/2

[4D (t − tc)]1/2

}[8]

where C�(tc) represents the grain boundary concentration atthe critical time tc. Figure 18 illustrates the thermo-inducednonequilibrium grain boundary segregation of P in 12Cr1MoVsteel.151,152 It is obvious that thermo-induced nonequilibriumgrain boundary segregation during isothermal annealing is char-acterized by a maximum grain boundary concentration.

Equation (5) indicates that the critical time becomes shorterif the temperature is elevated. At a rather low temperature, theduration time of a tensile test or the holding time prior to the

FIG. 18. Thermo-induced non-equilibrium grain-boundary seg-regation of P in 12Cr1MoV steel at 540 ◦C. (Reprinted withpermission from Li et al.151 Copyright 2002: Elsevier.)

test can be much shorter than the critical time. As a result, thediffusion of complexes towards the grain boundaries is domi-nant and the segregation amplitude of the impurities is far offits maximum (Figure 19a). When the temperature increases,diffusion of complexes and impurities are accelerated and con-sequently the critical time becomes shorter. When the durationor holding time equals the critical time at a given temperature,the grain boundary concentration of impurities reaches the max-imum (Figure 19b) and induces maximum embrittlement. Withfurther increase in temperature, the critical time can becomeeven shorter than the duration or holding time. Then, the reversediffusion of the impurity becomes dominant. In consequence,grain boundary segregation (Figure 19c) decreases and highductility is restored.

3.7. SummaryFrom the previous brief literature survey, it is seen that the

existing interpretations proposed for different alloys are quitedivers. In addition, researchers drew conflicting conclusionseven for a given alloy. For Hastelloy X, its intermediate tem-perature embrittlement was suggested to be caused by grainboundary shearing or sliding56 or thermo-induced nonequilib-rium grain boundary segregation of S.57,58 In the case of fillermetals, the embrittlement were ascribed to intergranular pre-cipitates92 or grain boundary sliding.95,96 The lack of consensusnecessitates a proper review of the preceding interpretationsto identify the most probable mechanism that yields a generalunderstanding of intermediate temperature embrittlement.

4. DISCUSSIONSIn Section 2.5, the generality of intermediate temperature

embrittlement of Ni and Ni-based superalloys has been con-firmed. Table 1 summarizes the preceding literature survey on

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FIG. 19. Schematic illustration of the interpretation due to thermo-induced non-equilibrium grain-boundary segregation.

Ni alloys that reveal intermediate temperature embrittlement. Itis obvious that each interpretation is accepted by different au-thors and supported by many experiments. This situation raisesdifficulties in the comprehension and control of intermediatetemperature embrittlement. In the following, existing interpre-tations and representative experiments are discussed in viewof the generality in order to develop a better understanding ofintermediate temperature embrittlement.

4.1. Discussion on the Interpretation due toIntergranular Precipitates

It is known that intergranular precipitates can reduce ductil-ity owing to stress and/or strain concentrations near precipitatesand thus cause easy initiation of a crack along grain boundary.Accordingly, the interpretation of intergranular precipitates isthe first and most popular for intermediate temperature embrit-tlement of Ni alloys.

In 1925, Merica and Waltenberg48 attributed the detrimentaleffects of S on the ductility of Ni to the formation of a Ni-Ni3S2 eutectic film at grain boundaries. On the contrary, Kraaiand Floreen88 considered that a liquid phase cannot be the solecause of embrittlement of Ni(S) alloys because low ductility isalready obtained at 427 ◦C (see Figure 9). Further observationsby transmission electron microscopy did not reveal any secondphase particles in grain boundaries of Ni(S) alloys.89,90 Besides,Thompson et al.167 concluded that a low melting grain boundaryphase Ni3S2 is not necessary for intergranular failure of Inconel718 doped with B, P, S, and/or C. Moreover, experimental re-sults of Brigham et al.168 indicate that Ni3S2 cannot form in Niat S levels of 20 ppm, while intermediate temperature embrit-

tlement is clearly observed at this and even lower S levels (seeFigure 9).88,91,93 Thus, intermediate temperature embrittlementof commercially pure Ni and Ni-based superalloys cannot bedue to the formation of Ni3S2, although this sulfides can giverise to embrittlement in principle.

In 1991, Vernot-Loier and Cortial34 believed that the minimain both reduction in area and impact strength of forged Alloy 625were induced by intergranular precipitates of M23C6 and M6Cat grain and twin boundaries. In 1994, Cortial et al.50 foundthe same minima in weld alloy 625 (see Figure 20) and relatedthe catastrophic effect on ductility and impact strength to theprecipitation of δ phase. In 2006, the decrease in ductility of In-conel 718 processed through powder metallurgy in combinationwith hot isostatic pressing route (see Figure 7) was attributed byAppa Rao et al.85 to decoration of prior particle boundaries withhighly stable and brittle oxides and MC carbides. In 2008, tensileductility minima of Ni-Cr filler alloys were suggested by Younget al.94 to correspond to partially or fully coherent second phaseprecipitates. They also related the ductility losses of InconelX-750, Inconel 718 and Monel K-500 to M23C6-type carbide,M6C-type carbides and γ ′ precipitate, respectively. However,these propositions cannot be applied to intermediate tempera-ture embrittlement in general. First, binary Ni(S),88–90 Ni(Bi)alloys (see Figure 21),30 and commercially pure Ni20092 werefound to be free of precipitates but extremely sensitive to in-termediate temperature embrittlement, which demonstrates thatintergranular precipitates are not necessary for embrittlement,especially not for the maximum. Second, in 1992, Miskovicet al.70 studied the intermediate temperature embrittlement ofthe as-cast and heat treated condition of IN 939 Ni-based su-peralloy. A ductility minimum appears at 800 ◦C for the as-cast

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TABLE 1Intermediate temperature embrittlement of Ni and Ni-based superalloys and corresponding interpretations

Alloy Manufacturing routes Hardening principleTemperature of ductility

minimumPresumed

interpretation

Alloy C-276 Wrought SS 800 ◦C32 DSAAlloy C-22 Wrought SS 700 ◦C54 DSAInconel 600 Wrought SS 760 ◦C59 GBSNi-Cr-W Wrought SS 600 ◦C41 GBSHastelloy X Wrought SS 649 ◦C56 GBS

500 ◦C57,58 TNGSWaspaloy Wrought AH 800 ◦C55 DSAUdimet 115 Wrought AH 850 ◦C36 DGPIN-738LC Wrought AH 650 ◦C33,53 DGPInconel 740 Wrought AH 750 ◦C52 —Nimonic 263 Wrought AH 800 ◦C52 —Alloy 625 Wrought AH 700 ◦C34,50,51 IPInconel X-750 Wrought AH 850 ◦C94 IPMonel K-500 Wrought AH 650 ◦C94 IPALLVAC 718PLUS Wrought AH 815 ◦C97 GPEInconel 718 Wrought AH 704 ◦C35 —

Wrought AH 760 ◦C94 —Cast (CC) AH 760 ◦C35 —

PM AH 850 ◦C85 IPIN939 Cast (CC) AH 800 ◦C70 —M963 Cast (CC) AH 800 ◦C69 DGPB1900 Cast (CC) AH 871 ◦C39 DGPM951 Cast (CC) AH 800 ◦C42 DGPMAR-M247 Cast (CC) AH 760 ◦C40 DGPCM247LC Cast (CC) AH 650 ◦C37,38 DGP

Cast (DS) AH 760 ◦C37,38 DGPGTD-111 Cast (DS) AH 750 ◦C71–74 DGPPWA 664 Cast (DS) AH 815 ◦C75 —PWA 659 Cast (DS) AH 815 ◦C75 —DZ951 Cast (DS) AH 760 ◦C76,77 DGPPW 1484 Cast (SC) AH 649 ◦C78 —DD3 Cast (SC) AH 800 ◦C79 —APK-1 PM AH 760 ◦C86 GPENi270 — — 800 ◦C91,98 GPENi(Bi) alloy — — 750 ◦C30 TNGSNi(S) alloy — — 649 ◦C88–91 NGSFe-36Ni — — 950 ◦C99 EGSCP-Ni — — 700 ◦C31 GBS

— — 760 ◦C87 —FM52 — — 860 ◦C94 IP

— — 750 ◦C92,93,95,96 GBS

CP-Ni: commercially pure Ni; CC: conventionally cast; DS: directional solidification; SC: single crystal; PM: powder metallurgy; SS: solutionstrengthened; AH: age hardened; DSA: dynamic strain aging; GBS: grain boundary shearing or sliding; TNGS: thermo-induced nonequilibriumgrain boundary segregation; DGP: decohesion of glide plane; IP: intergranular precipitates; GPE: gas phase embrittlement; EGS: equilibriumgrain boundary segregation.

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FIG. 20. (a) Ductility and (b) impact strength of alloy 625 vs.temperatures after data from Cortial et al.50

condition while no intermediate temperature embrittlement wasfound for the fully heat treated condition (solution treatment andaging). In 2004, He et al.69 observed that a ductility minimum ofM963 superalloy locates at 800 ◦C under solution treatment anddisappears under age treatment. In the same year, the minimumductility of IN-738LC superalloy in standard heat treatmentcondition was found by Sharghi-Moshtaghin and Asgari33 tolocate at 650 ◦C. However, no drop in ductility was observedfor samples aged for rather long time (1500 and 3000 hours).In 2007, Xia et al.76 found that the DZ951 superalloy showsevident intermediate temperature embrittlement after standardheat treatment (see Figure 22) while embrittlement vanishesas the isothermal exposure time extends (100, 500, 1000 and2000 hours). In 2009, the tensile properties of 720Li superalloywere measured by Gopinath et al.169 No intermediate tempera-ture embrittlement of this alloy after solution treatment and twostages of long time aging (see Figure 23) was observed. If the in-terpretation of intergranular precipitates were valid, these super-alloys should exhibit maximum embrittlement after long-timeexposure or aging as well. This contradicts to the experimentalresults and allows the conclusion that intergranular precipitatesare not a precondition for intermediate temperature embrittle-ment. Third, single crystal Ni-based superalloys, such as PW1484,78 DD3,79 CMSX-4,80,81 and CMSX-10,82 were provento have evident intermediate temperature embrittlement. Obvi-ously, the embrittlement of single crystal Ni-based superalloyscannot be explained by intergranular precipitates just becauseof the absence of grain boundaries.

4.2. Criticism on the Interpretation due to GrainBoundary Shearing or Sliding

Grain boundary shearing or sliding is accepted widely as themechanism of creep embrittlement. The interpretation of grainboundary shearing or sliding was first proposed by Rhines andWray31 in 1961 for intermediate temperature embrittlement. In1970, Shapiro and Dieter59 explained the embrittlement of hotrolled Inconel 600 (see Figure 24) in virtue of this idea. In 2004,the intermediate temperature embrittlement of two Ni-basedfiller metals (FM-52 and FM-82) were studied by Ramirez andLippold95,96 and also postulated to be the result of a creep-likemechanism associated with grain boundary sliding.

Although the embrittlement caused by grain boundary shear-ing or sliding was confirmed by many experiments, this inter-pretation is insufficient to be a general effect. First, if the in-termediate temperature embrittlement of binary Ni(S)88–90 andNi(Bi)30 alloys as well as commercially pure Ni20092,93 andimpurity added Ni87 resulted from grain boundary shearing orsliding, highly pure Ni must exhibit a maximum in embrittle-ment likewise, since grain boundary shearing or sliding willalso work in pure metal. Whereas, experimental observationsclearly exclude intermediate temperature embrittlement in thecase of highly pure Ni (see Figures 8 to 10).30,87,88 Second, in-termediate temperature embrittlement of single crystal Ni-based

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FIG. 21. Optical microscopy images of samples after tension at (a) 400 ◦C, (b) 750 ◦C, and (c) 825 ◦C as well as transmissionelectron microscopy image of samples heat treated at (d) 750 ◦C for 45 min. (Reprinted with permission from Zheng et al.30

Copyright 2011: Elsevier.) (Color figure available online).

superalloys78–82 cannot be explained by grain boundary shearingor sliding. Third, the interpretation ascribed ductility recoveryto recrystallization at high temperatures. This contradicts tothe fact that the ductility recovery was observed in alloys thatdefinitely show no recrystallization at temperatures above theductility minimum. For example, in 1972, Arkoosh and Fioer56

demonstrated that Hastelloy X (see Figure 3), which exhibitsminimum ductility at 650 ◦C, does not recrystallize even at760 ◦C. In 2011, Zheng et al.30 showed that a Ni(Bi) alloy re-veals a minimum ductility between 700 ◦C and 750 ◦C while no

recrystallization was observed in the whole studied temperaturerange (see Figure 21).

Although they did not support the idea that ductility recoveryis due to recrystallization, Arkoosh and Fioer56 still linked theductility minimum of Hastelloy X to grain boundary shearingin combination with fine secondary carbide precipitates at grainboundaries. They supposed that above the temperature of theductility minimum, grain boundary shear can still occur, butthe matrix is depleted of strengthening elements by precipita-tion. So, ductility is restored by softening the matrix. However,

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FIG. 22. Influence of thermal exposure on the tensile elongationof DZ951 alloy.76 SHT: standard heat treatment. (Reprinted withpermission from Xia et al.76 Copyright 2007: Elsevier.) (Colorfigure available online).

this idea combining grain boundary shearing with precipitatesstill does not fit the intermediate temperature embrittlement ofNi(S) alloys,88–90 commercially pure Ni200,92,93 and Ni(Bi) al-loy30 that do not contain any precipitates at all. Moreover, thisidea cannot explain the disappearance of embrittlement of IN-738LC,33 M963,69 IN939,70 DZ951,76 and 720Li169 superalloysafter long time of exposure or aging since precipitates are stillpresent after long isothermal treatments.

FIG. 23. Tensile ductility of 720Li superalloy at different tem-peratures. (Reprinted with permission from Gopinath et al.169

Copyright 2009: Elsevier.) (Color figure available online).

FIG. 24. Total shear strain vs. temperature of Inconel 600 afterdata from Shapiro and Dieter.59

4.3. Remark on the Interpretation due to Gas PhaseEmbrittlement

In 1981, Bricknell and Woodford91 studied the tensile duc-tility of commercially pure Ni270. Specimens were first ex-posed at 1000 ◦C for 200 hours to air or to vacuum (1.33×10−3

Pa) prior to testing. Then, they were furnace cooled and ten-sile tested in highly pure argon at temperatures from 25 ◦Cto 1000 ◦C. The reduction in area indicates maximum em-brittlement around 800 ◦C. Evidently, the embrittlement afterexposure in air is more significant than that in vacuum (seeFigure 25). Further investigation identified oxygen penetrationalong grain boundaries as the reason for embrittlement. It wasdemonstrated that oxygen can have considerable effect, evenif specimens are exposed to very low partial pressure (<1.33× 10−5 Pa). Therefore, Bricknell and Woodford91 ascribed theductility loss to grain boundary pinning by oxygen penetrationand ductility recovery at higher temperatures to dynamic recrys-tallization. Even though this idea seems to explain intermediatetemperature embrittlement successfully, it is unsatisfactory inview of the entire set of experiments. First, if very low partialpressure of oxygen (<1.33×10−5 Pa) was sufficient to causeembrittlement and dynamic recrystallization was capable to re-store ductility,91 one had to conclude that even highly pure Nishould display maximum embrittlement when normally tested.But this deduction is clearly in contrast with the fact that highlypure Ni does not embrittle.30,87,88 Second, the heat treatmentsof Ni270 before tensile testing by Bricknell and Woodford91

differ from those of binary Ni alloys significantly. The binary

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FIG. 25. Intermediate temperature embrittlement of Ni270.(Reprinted with permission from Bricknell and Woodford91

Copyright 1981: Springer Science + Business Media.)

Ni(S)88 and Ni(Bi)30 alloys were held for only 45 minutes atthe respective testing temperatures (427 ◦C to 760 ◦C for Ni(S)alloy and 400 ◦C to 825 ◦C for Ni(Bi) alloy, respectively), whileBricknell and Woodford91 subjected the Ni270 samples to muchlonger time and higher temperature heat treatment (200 hours at1000 ◦C). This treatment made oxygen produce a pronouncedeffect on pinning grain boundaries in subsequent tensile testing,but misled, to an invalid interpretation for the general cause ofintermediate temperature embrittlement.

Similar arguments were presented for particle-strengthenedNi-based superalloys. In 1983, Prakash et al.86 suggested thatgrain boundaries of a powder metallurgy APK-1 alloy are weak-ened by oxygen adsorption and therefore cracks are initiatedat stress concentrations ahead of dislocation pile-ups at inter-mediate temperatures. A large volume fraction of fine γ ′ pre-cipitate promoted a tensile ductility minimum, since environ-mental oxidation affects the bond between particle and matrixadversely. In 2008, a series of uniaxial tensile tests of ALL-VAC 718PLUS superalloy were carried out by Hayes.97 It wassupposed that grain boundary cracking is due to the penetrationof oxygen along grain boundaries and subsequent oxidation ofgrain boundary particles (carbides and possibly delta phase).However, these interpretations are unsatisfactory for intermedi-ate temperature embrittlement of age-hardened superalloys ingeneral. First, they cannot explain the frequently observed dis-appearance of intermediate temperature embrittlement after pro-longed aging of alloys that contain both fine γ ′ precipitate andcarbides.33,69,70,76,169 Second, Blankenship and Henry170 con-cluded that environmental embrittlement of Rene88 DT and In-

FIG. 26. Ductility (characterized by elongation at fracture, Ef)minima in Udimet 115 superalloy. (Reprinted with permissionfrom Jensen and Tien36 Copyright 1985: Springer Science +Business Media.)

conel 718 can be ruled out due to the fact that the embrittlementobtained in vacuum was the same as that in air. This conclusionstrongly suggests the ideas of Prakash et al.86 and Hayes97 tobe invalid since both alloys have a large volume fraction of γ ′

precipitates and grain boundary particles.Accordingly, gas phase embrittlement is inadequate as a gen-

eral interpretation of intermediate temperature embrittlement.At the same time, one must clearly state that gas phases in-duce embrittlement definitely although it is not the generalmechanism of ITE. For example, dynamic embrittlement wasproven as a generic type of brittle intergranular fracture inwhich the crack propagation is controlled by grain boundarydiffusion of an embrittling element. This element may comefrom the material itself or from the surrounding atmosphere(oxygen in particular).10,100,101,171–174 This phenomenon was ex-perimentally proven in Ni-based superalloys10,100,101,171–174 andtheoretically supported by Bassani et al.,175,176 Vitek,177 Kruppet al.,10,100,101,173 and McMahon et al..10,100,172,173,174,178

4.4. Evaluation of the Interpretation due to Glide PlaneDecohesion

In 1985, Jensen and Tien36 investigated the temperatureand strain rate dependence of stress-strain behavior in theγ ′ precipitation-strengthened Udimet 115 Ni-based superal-loy. The elongation is shown in Figure 26 as a function of

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200 L. ZHENG ET AL.

temperature and strain rate. The curves demonstrate an evidentductility minimum in the vicinity of 850 ◦C. The appearanceof the fracture surface at the ductility minimum is concluded tobe glide plane decohesion at the γ /γ ′ interface. In 1995, Bettgeet al.53 investigated the tensile properties of IN-738LC superal-loy and explained the intermediate temperature embrittlementin terms of the same concept. In 2004, Kim et al.,39 Sajjadiet al.,71–74 and He et al.69 also used this idea to explain the lowductility of the cast B1900, GTD-111, and M963 superalloys, re-spectively. In 2008, ductility minima in conventionally cast anddirectional solidification CM247LC superalloy were consideredby Kim et al.37,38 as resulting from the localized deformation atthe γ /γ ′ interface and the consequent glide plane decohesion.

Definitely, glide plane decohesion will result in fast proroga-tion of cracks along a certain plane and subsequent low ductilityof Ni alloys. However, as the decisive reason for intermediatetemperature embrittlement, this interpretation is unsatisfactory.First, intermediate temperature embrittlement is mainly reportedin conjunction with intergranular fracture,32,49,57,58,91,97–99 whichcontradicts a possible mechanism of glide plane decohesionthat would result in transgranular fracture. Second, commer-cially pure and impurity added Ni30,87,88,92,93 as well as solutionstrengthened Ni-based superalloys32,41,54,56–58 reveal no γ /γ ’interfaces but show the embrittlement clearly. Third, this inter-pretation cannot be responsible for the disappearance of embrit-tlement after long time exposure or aging.33,69,70,76,169

4.5. Judgment of the Interpretation due to DynamicStrain Aging

It is accepted widely and has been frequently demonstratedthat dynamic strain aging can give rise to the embrittlement ofalloys, since diffusing solute atoms will interact with mobiledislocations and thereby impede the movement of dislocations.

In 2008, Roy et al.32,54,55 studied the tensile properties of Al-loy C-22, Alloy C-276 and Waspaloy from ambient temperatureto 1000 ◦C. They found that these alloys reveal clear intermedi-ate temperature embrittlement and exhibit serrations in the en-gineering stress-strain diagrams (see Figures 27 and 28). Sincedynamic strain aging occurs within a susceptible temperatureregime similar to that of tensile ductility loss, the embrittlementwas related to dynamic strain aging.32,54,55 This interpretationseems to be particularly attractive, because it could explain theintermediate temperature embrittlement of not only Ni-basedsuperalloys but also binary Ni alloys as well as impurity addedNi. Since the dynamic strain aging originates from the interac-tion of dislocations with solute atoms, it also naturally explainsthe disappearance of intermediate temperature embrittlement ofhighly pure Ni.

Nevertheless, this interpretation does not yield a compre-hensive description of intermediate temperature embrittlement.First, ductility minimum of Alloy C-22 appears at 700 ◦C(see Figure 27a) but stress-strain curves show the most obvi-

FIG. 27. (a) Tensile ductility and (b) stress-strain curves vs.temperature of Alloy C-22. (Reprinted with permission fromRoy et al.54 Copyright 2008: Elsevier.) (Color figure availableonline).

ous behavior of dynamic strain aging at 600 ◦C (see Figure27b).54 Ductility minimum of Alloy C-276 appears at 800 ◦C(see Figure 28a) while dynamic strain aging in stress-straincurve has already disappeared at this temperature (see Fig-ure 28b).32 Ductility minimum of Waspaloy appears at 800 ◦Cwhereas the stress-strain curves display clear dynamic strainaging only within the range of 300 ◦C to 600 ◦C.55 Similarly,Jensen and Tien36 confirmed that a ductility minimum of Udimet115 superalloy exists in the vicinity of 850 ◦C (see Figure 26)while the serrated flow can only be found at temperatures from538 ◦C to 649 ◦C. These comparisons indicate that dynamicstrain aging and intermediate temperature embrittlement appearwithin distinct temperature ranges. Second, Shapiro and Dieter59

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FIG. 28. (a) Tensile ductility and (b) stress-strain curves vs.temperature of Alloy C-276. (Reprinted with permission fromRoy et al.32 Copyright 2008: Elsevier.) (Color figure availableonline).

assured that a ductility minimum of Inconel 600 exists in thevicinity of 760 ◦C (see Figure 24) while there is no occur-rence of serrated flow in torque-twist tests (see Figure 29). Thesame situation was observed in Fe-36Ni alloy by Ben Mostefaet al.99 and in Ni(Bi) alloy by Zheng et al.30 On the other hand,Gopinath et al.169 proved that intermediate temperature embrit-tlement does not occur in 720Li superalloy (see Figure 23) butdynamic strain aging appears (see Figure 30). All these resultsdemonstrate that intermediate temperature embrittlement anddynamic strain aging do not appear concurrently. Third, thefrequent intergranular fracture within the temperature range ofintermediate temperature embrittlement32,49,57,58,91,97–99 cannotbe explained by dynamic strain aging, since dynamic strain ag-ing would cause transgranular fracture mainly. Fourth, it should

FIG. 29. Torque-twist curves of Inconel 600.59

be noted that a minimum of impact strength can also occur incombination with intermediate temperature embrittlement,34,50

as illustrated in Figure 20. Consequently, one can expect thatintermediate temperature embrittlement and impact trough area consequence of the identical microscopic mechanism. As weknow, dynamic strain aging is the result of the diffusion of so-lute atoms in matrix and the accumulation of these solute atomsaround dislocations. Therefore, dynamic strain aging will not beeffective during impact since there is not enough time for soluteatoms to diffuse and accumulate.

FIG. 30. Strain-stress curve of Ni-based superalloy 720Li.(Reprinted with permission from Gopinath et al.169 Copyright2009: Elsevier.)

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202 L. ZHENG ET AL.

FIG. 31. Calculated grain boundary segregation of Bi at equi-librium state (solid) and after 45 min heat treatment at differenttemperatures (dashed, dotted lines). (Reprinted with permissionfrom Zheng et al.30 Copyright 2011: Elsevier.) (Color figureavailable online).

4.6. Comment on the Interpretation due to GrainBoundary Segregation

4.6.1. Equilibrium Grain Boundary SegregationAs described in section 3.6.1, equilibrium grain boundary

segregation could in principle produce intermediate tempera-ture embrittlement. In 2011, thermodynamics and kinetics ofgrain boundary segregation of Bi in Ni(Bi) alloy were calcu-lated by Zheng et al.30 (see Figure 31) based on data of thesegregation energy of Bi in Ni129 and the averaged diffusion co-efficients of different impurities in Ni.179 In order to predict themaximum grain boundary segregation of Bi at 700 ◦C (Group I),725 ◦C (Group II), or 750 ◦C (Group III) in accordance with theexperimental observations (see Figure 10), the necessary valuesof pre-exponential factors of Bi diffusion are far off realisticranges (two orders of magnitude higher at least). Obviously,equilibrium grain boundary segregation theory is not applica-ble to describe the intermediate temperature embrittlement ofNi(Bi) alloy in a quantitative manner.

In 1988, Ben Mostefa et al.99 attributed the ductility drop ofFe-36Ni alloys to grain boundary segregation of S and the ductil-ity recovery to dynamic recrystallization. However, no dynamicrecrystallization of Ni(Bi) alloy during tensile test is observed asshown in Figure 21.30 So, the interpretation combining equilib-rium grain boundary segregation with dynamic recrystallizationis not an acceptable alternative.

4.6.2. Nonequilibrium Grain Boundary SegregationIn 1964, the tensile ductility of Ni containing S of 1 to 50 ppm

at 427 ◦C to 760 ◦C was studied by Kraai and Floreen,88 as shown

in Figure 9. The results revealed that a small amount of S (5 to20 ppm) produces intermediate temperature embrittlement. In1969, Floreen and Westbrook89 measured the microhardness ofthese alloys and concluded that the local increase in micro-hardness near grain boundaries is induced by grain boundarysegregation of S. In 1974, Westbrook and Floreen90 re-analyzedthe grain boundary microhardness and confirmed that the kinet-ics of grain boundary hardening is consistent with a vacancyassisted, thermol-induced nonequilibrium grain boundary seg-regation of S. In their review of the effects of harmful and ben-eficial elements on the properties of Ni-based superalloys, Holtand Wallace180 pointed out that thermo-induced nonequilibriumgrain boundary segregation is a more satisfactory model to ex-plain the grain boundary hardening in Ni(S) alloys than eitherthe model of precipitations or that of equilibrium grain bound-ary segregation involving S without formation of second phaseparticles. Although Floreen et al.88–90 determined the interme-diate temperature embrittlement of Ni(S) alloys and proved thegrain boundary segregation of S to be of nonequilibrium type,they did not correlate the embrittlement to the nonequilibriumgrain boundary segregation of S.

It is well known that S segregates strongly at grain bound-aries of Ni and Ni-based superalloys and induces severe embrit-tlement.57,58,120,180–185 Recently, Wang et al.57,58 studied the in-termediate temperature embrittlement of Hastelloy X as shownin Figure 32. It is clear that the maximum grain boundary seg-regation of S corresponds to the maximum embrittlement of thealloy. Moreover, the grain boundary segregation of S was con-firmed to be controlled by thermo-induced nonequilibrium grainboundary segregation in view of a clear maximum concentra-tion amplitude. Thus, Wang et al.57,58 proved that intermediatetemperature embrittlement of Hastelloy X is induced by thermo-induced nonequilibrium grain boundary segregation of S. In2011, Zheng et al.30 calculated the critical time of Bi in Ni(Bi)alloy based on summarized diffusion coefficients of impuritiesin Ni (see Figure 33). Remarkably, the temperatures correspond-ing to the critical time of 45 min, which was the holding timeof samples prior to tensile test in corresponding experiments,fall to good approximation in the temperature range of 700 ◦Cto 750 ◦C in which the maximum embrittlement of Ni(Bi) alloywas observed (see Figure 10). Therefore, the study confirmsquantitatively the mechanism of thermo-induced nonequilib-rium grain boundary segregation for intermediate temperatureembrittlement.

4.6.2.1. Arguments in favor of nonequilibrium grainboundary segregation theory. From the sketched model147–150

and above discussions with respect to thermo-induced nonequi-librium grain boundary segregation, it is clear that this inter-pretation is applicable to describe the intermediate tempera-ture embrittlement of binary Ni alloys, impurity added Ni,and Ni-based superalloys. In agreement with the stated ex-periments, it can naturally explain the disappearance of in-termediate temperature embrittlement in the case of highlypure Ni.

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FIG. 32. (a) Reduction in area and (b) grain boundary S segre-gation vs. temperature for Hastelloy X.57,58

Furthermore, the thermo-induced nonequilibrium grainboundary segregation theory had been successfully applied toexplain the reversible temper embrittlement characterized byimpact strength in steels.186,187 So, this theory can be used toexplain the concomitant minimum of impact strength of alloy625 observed by Vernot-Loier and Cortial34 and Cortial et al.50

as well.In particular, intermediate temperature embrittlement of

IN939,70 M963,69 IN-738LC,33 DZ951,76 and 720Li169 super-alloys disappear after sufficient time of exposure or aging.These observations are difficult to be rationalized by the alterna-tive interpretations but easy to be predicted by thermo-inducednonequilibrium grain boundary segregation.147–150 As describedin the model, nonequilibrium segregation level disappears anddistribution of species reaches equilibrium as time approaches

infinity. The disappearance of segregation naturally leads to thatof intermediate temperature embrittlement during subsequenttensile tests.

It should be mentioned that many kinds of solute atoms (S,P, Pb, Bi, B, Mo, W, Cr, Nb, Fe, et al.) and different types ofcarbides (MC, M6C, M23C6, M7C3, et al.) as well as topolog-ical phases (δ, Laves, σ , µ, η, et al.) may exist in Ni-basedsuperalloys. Therefore, any exact explanation of intermediatetemperature embrittlement in a certain alloy needs to take intoaccount the evolution of second phases, and the interactionsbetween different kinds of solute atoms117 or between soluteatoms and precipitating phases.137,138 But as a general guideline, thermo-induced nonequilibrium grain boundary segrega-tion is the most probable concept to understand the ductilitydrop of Ni alloys in intermediate temperature range.

4.6.2.2. Extended application of nonequilibrium grainboundary segregation theory. Although the interpretation ofthermo-induced nonequilibrium grain boundary segregation canexplain most experiments about intermediate temperature em-brittlement, the present model is still not perfect. Single crystalsuperalloy PW 1484 (see Figure 6),78 DD3,79 CMSX-4,80,81 andCMSX-1082 were proven to display intermediate temperatureembrittlement. Since these alloys have not grain boundary, theinterpretation of thermo-induced nonequilibrium grain bound-ary segregation is not applicable evidently. In 1997, Liu et al.188

investigated the effects of several impurities on the γ /γ ′ inter-face cohesion by employing first-principles electronic structurecalculations. They found that S is the most deleterious speciesto reduce the γ /γ ′ interface cohesion. In 2003, Chen et al.189

suggested based on calculation by first principle approach that Swill segregate to the γ /γ ′ interface and embrittle single crystalsuperalloys. Moreover, Dong et al.117 determined the S con-centration at the carbide/matrix interface in Inconel 718 su-peralloy to be 4.74 wt%, which represents approximately onethousand times the bulk concentration. They also found that Sprefers to segregate at carbide/matrix interface rather than atgrain boundary. This phenomenon was confirmed by Nettle-ship and Wild125 and Ben Mostefa et al.99 In view of these facts,the intermediate temperature embrittlement of single crystal Ni-based superalloy probably stems from nonequilibrium “phaseinterface” rather than grain boundary segregation of S or otherharmful elements. For certain, this presumption still needs tobe identified although it seems reasonable. If it can be con-firmed, nonequilibrium “interface” (including grain boundaryand phase interface) segregation will be the general mechanismof intermediate temperature embrittlement of Ni and Ni-basedsuperalloys. For the analytical description of the segregation toan arbitrary interface, the thermodynamic and kinetic equationsof grain boundary segregation can be applied since both typesof segregation root certainly in the decrease in free energy of thewhole system by the diffusion of solute atoms. As for the exper-imental investigation, many techniques have been developed todetect grain boundary segregation, such as particle tracking au-toradiography,190 secondary ion mass spectroscopy,119,135,136,191

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204 L. ZHENG ET AL.

FIG. 33. Calculated critical time of grain boundary segregation of Bi at different temperatures for a probable range of Bi diffusivitiesin the Ni(Bi) binary alloy.30 (Reprinted with permission from Zheng et al.30 Copyright 2011: Elsevier.) (Color figure availableonline).

X-ray photoelectron spectroscopy,135,136,191,192 Auger elec-tron spectroscopy,114,135,136,191 high resolution electron mi-croscopy,135,191,193 electron energy loss spectroscopy,191,194 andatom probe tomography.135,136,191,195–198 For future composi-tional analysis of interphase boundaries, the most preferablemethod is atom probe tomography, since it allows a three-dimensional reconstruction of the spatial distribution of var-ious atomic species with single atom sensitivity and a spa-tial resolution in the range of 2–4 Å.199–204 Furthermore,the enrichment of Ti, Al, Mg, Si, P, and Re at γ /γ ′ inter-face in Ni-based superalloys were already detected by thistechnique.205–208 So, the quantitative detection of impuritysegregation at interphase boundaries can be expected and theirimpact on intermediate temperature embrittlement can be clari-fied consequently.

In 1971, Dix and Savage209 found that increasing the de-formation rate of an age-hardened alloy from 6.77×10−5 to6.77×10−3 m·s−1 decreases the depth of the elongation mini-mum from 10% to 30%. In 2004, this phenomenon was provenagain in solution strengthened alloy 625 by Shankar et al.51 Inaddition, Jensen and Tien36 pointed out in 1985 that the ductilityminimum in Udimet 115 superalloys shifts to lower tempera-tures at lower strain rates (see Figure 26). The ductility minimumin the Al-5.2%Mg alloy was also found to shift from 380 ◦Cat the highest strain rate (10−1 s−1) to 230 ◦C at the sloweststrain rate (10−5 s−1).210 Clearly, the outlined theory of thermo-induced nonequilibrium grain boundary segregation cannot ex-plain the influence of strain rate on intermediate temperatureembrittlement easily. However, it is known that excess vacan-cies can be created during plastic deformation in metals211 and

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form vacancy-impurity complexes212 afterwards. This is pre-requisite for nonequilibrium grain boundary segregation.147–150

In 2008, Chen et al.213 explained the strain rate dependence ofintermediate temperature embrittlement in the Al-5.2% Mg al-loy by a deformation-induced nonequilibrium grain boundarysegregation of Na. This provides an approach to investigate theinfluence of strain rate. On the other hand, in 2009, Xu et al.214

re-explained the shift in the temperature of ductility minimum inthe discussed Al-5.2%Mg alloy. They pointed out that decreas-ing strain rate in tensile tests at elevated temperature meansextending the isothermal testing time and thereby shifting thetemperature of significant embrittlement to a lower one. In 2011,Zheng et al.30 proved that the time to tensile fracture of Ni(Bi)alloy at 700 ◦C and 750 ◦C, at which the ductility minimumappears, takes only around 30 seconds (at a strain rate of 10−2

s−1). In such a short time, deformation cannot cause evidentgrain boundary segregation of Bi and subsequent embrittlementof Ni(Bi) alloy. So, deformation-induced nonequilibrium grainboundary segregation is not the primary reason for intermediatetemperature embrittlement in tensile test of high strain rate butit should be taken into account if strain rate is slow.

In a word, the nonequilibrium interface segregation theory,including the effect of nonequilibrium vacancies that are eitherinduced by quenching or by plastic deformation, and involvingthe effect of strain rate on segregating behavior of impurities,can probably describe the nature of intermediate temperatureembrittlement in a consistent way.

4.7. SummaryIn 1975, Fiore49 concluded that different alloy systems may

exhibit ductility minima for different reasons. In fact, studiesconcerning highly pure Ni,30,87,88 commercially pure Ni,92,93

impurity added Ni,87 Ni(S) alloys,88 and Ni(Bi) alloy30 had al-ready indicated explicitly that intermediate temperature embrit-tlement is an impurity effect and that impurity levels of severalppm are sufficient to create such an embrittlement. Thus, theinterpretations due to grain boundary shearing or sliding, gasphase embrittlement, and glide plane decohesion are not sat-isfactory in general. Intergranular precipitates is a kind of im-purity effect. However, commercially pure and impurity addedNi without precipitates also revealed intermediate temperatureembrittlement. Thus, the effect of intergranular precipitates canbe excluded as the main cause. Although dynamic strain agingresults from impurity effect, the interpretation of dynamic strainaging is nevertheless not applicable since the temperature rangeof intermediate embrittlement does not properly correlate withthat of dynamic strain aging.

The remaining, quite promising interpretation of thermo-induced nonequilibrium grain boundary segregation is not ca-pable to explain the intermediate temperature embrittlementof single crystal superalloy and the influence of strain rate atpresent. So, an entirely satisfactory interpretation of interme-diate temperature embrittlement has not been derived yet, al-

though a considerable amount of work had already been done.A “nonequilibrium interface segregation” taking into accountthe additional effect of strain rate is probably the most generalmechanism. It is well known that interfacial segregation is aninteraction between structural defects and impurities. In slightgeneralization, vacancy-impurity complexes produced duringtensile deformation may also be regarded as an interaction of im-purities with “zero-dimensional” structural defects. Thus, froma more general point of view, the complete mechanism of in-termediate temperature embrittlement should be rationalized asan interaction mechanism of impurities with a type of structuraldefects. Further work is clearly required at this point.

5. FUTURE DIRECTIONSIn this section, future investigations on the mechanism of

intermediate temperature embrittlement are suggested in virtueof the preceding review.

First, most reports of intermediate temperature embrittlementshare the common feature that the fracture mode is transgranu-lar outside the ductility minimum temperature range and inter-granular within this range.32,49,57,58,91,97–99 Such fractography ofAlloy C-276 is exemplified in Figure 34.32 However, a fewexperiments manifested entirely antithetical fracture modes.Jensen and Tien,36 Arkoosh and Fiore56 as well as Sharghi-Moshtaghin and Asgari33 reported that the fracture appearancesof Udimet 115, Hastelloy X and IN-738LC are opposite to thegeneral trend (see Figure 35). The results of tensile ductilityand fractography of a Ni-Cr superalloy obtained by the presentauthors are illustrated in Figure 36 and Figure 37. The inter-mediate temperature embrittlement are evident but the percent-age of intergranular surfaces exhibits a decreasing trend withincrease in tensile temperature, which differs from both thecommon and the opposite fractography. In 1985, Kandra andCosandey215 demonstrated that the ductility of a Ni-20Cr alloyis low and the fracture mode is 100% intergranular at the tensiletemperature of 700 ◦C. As the Ce content in this alloy increases,the ductility increases and the fracture path reverts from inter-granular to transgranular. In 1988, the fractography of GH33and Inconel 718 superalloys were found to change from inter-granular to transgranular fracture mode by adding minor amountof Mg.216,217 In 2008, Yan et al.218 found that the tensile frac-ture mode of a directionally solidified superalloy transfers fromintergranular to transgranular as increasing the B content at870 ◦C. After summarizing the chemical compositions of thosealloys with antithetical fracture modes, it was found that all ofthem contain beneficial elements (B, Zr, and Mg). It is wellknown that these elements can enhance grain boundary cohe-sion of Ni alloys significantly by themselves and/or interac-tions with impurities.124,180,216–229 Doping of these elements (aswell as rare earth, such as Ce215,230 and Hf180,229,231) certainlymakes the fractography complicated, especially when elementsare doped together. Therefore, further work with emphasis on the

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206 L. ZHENG ET AL.

FIG. 34. Fractographs of Alloy C-276. Predominantly ductilefailures characterized by dimpled microstructures can be seenwhen the sample was tested at (a) 600 ◦C while intergranularbrittle failures at (b) 700 ◦C. (Reprinted with permission fromRoy et al.32 Copyright 2008: Elsevier.)

particular role of beneficial elements should be initiated to dis-cover the reason of abnormal fracture of specific superalloys.

Second, the influence of the state of applied stress on the in-termediate temperature embrittlement should be investigated. In1970, the intermediate temperature embrittlement of hot rolledInconel 600 was observed by Shapiro and Dieter59 using a hottorsion tester (see Figure 24). Figure 24 conveys that intermedi-ate temperature embrittlement can be produced not only by ten-sile stress but also by shear stress, and implies that intermediatetemperature embrittlement is independent of the type of appliedstress. So, it is expected that intermediate temperature embrit-tlement will also appear under the application of compressivestress, whereas no such an investigation exists in the literature so

FIG. 35. Fractographs of Hastelloy X. (a) Intergranular fracturemode outside and (b) crystallographic-type facets within theductility minimum temperature range. (Reprinted with permis-sion from Jensen and Tien36 Copyright 1985: Springer Science+ Business Media.)

far. It is well known that different types of applied stresses leadto various deformation behaviors of materials and consequentlydiverse embrittling degrees. Since superalloys normally oper-ate under a complex stress state, the influences of each stressand their coupling on intermediate temperature embrittlementshould be clarified. Rod parameter µσ , soft coefficient αS, andstress triaxiaty Rσ are commonly-used stress state parameters instudying the deformation and fracture of metal materials. Theyare defined as232–234

µσ = 2σ2 − σ1 − σ3

σ1 − σ3[9]

αS = 1

2

σ1 − σ3

σ1 − ν (σ2 + σ3)[10]

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FIG. 36. Tensile ductility of a Ni-Cr superalloy characterized byreduction in area and elongation. (Color figure available online).

Rσ =√

2 (σ1 + σ2 + σ3)

3√

(σ1 − σ2)2 + (σ2 − σ3)2 + (σ3 − σ1)2[11]

where σ 1 > σ 2 > σ 3 are principal stresses. Since these pa-rameters can describe any complex stress state, the degree ofintermediate temperature embrittlement may be correlated tothem in further experiments.

Third, the relationship between intermediate temperature em-brittlement and strain rate should be ascertained. It was alreadyconfirmed that the width, depth and position of intermediatetemperature embrittlement will vary with strain rate. Furtherwork should be carried out to reveal the details of this varia-tion. Besides, dynamic strain aging may appear when differentstrain rates are applied. Its additional influence should also beclarified.

Fourth, the theory of thermo-induced nonequilibrium grainboundary segregation needs to be further developed. Atpresent, this theory cannot explain the intermediate temperature

FIG. 37. Fractographs of a Ni-Cr superalloy at (a) ambient temperature, (b) 400 ◦C, (c) 800 ◦C and (d) 900 ◦C, showing adecreasing trend in the percentage of intergranular surfaces with increase in tensile temperature.

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208 L. ZHENG ET AL.

embrittlement of single crystal superalloy and the influence ofstrain rate. Hence, it should be developed into “nonequilibriuminterface segregation” with taking into account the influence ofstrain rate on the segregation of impurities. Definitely, the con-cept of nonequilibrium interface segregation should be validatedby measuring the compositions at interphase boundaries apply-ing modern microscopic analysis tools, for example, atom probetomography. Moreover, the correlation of nonequilibrium inter-face segregation with its impact on deformation properties maybe a suggested way to characterize intermediate temperatureembrittlement when the strain rate is slow.

These few conceptual suggestions indicate that extended,further experimental work is required to reveal the factors thatinfluence intermediate temperature embrittlement and clarifythe nature of this fundamental effect in Ni and Ni-based super-alloys.

6. CONCLUSIONFrom the survey of experimental results, it is clear that in-

termediate temperature embrittlement is a general feature of Niand Ni-based superalloys.

It is worked out that rather different interpretations wereproposed for intermediate temperature embrittlement up tonow, i.e., intergranular precipitates, grain boundary shearingor sliding, gas phase embrittlement, decohesion of glide plane,dynamic strain aging, and grain boundary segregation ofimpurities.

According to the detailed discussions regarding these inter-pretations, it is concluded that at present no satisfactory descrip-tion exists for a complete understanding of intermediate tem-perature embrittlement. From the presented systematic evalua-tion, the mechanism of “nonequilibrium interface segregation”,however, may be the one to provide a generally accepted com-prehension.

To perform a more entire analysis of intermediate tempera-ture embrittlement, further work is required with respect to theabnormal fracture appearances of specific alloys, the influenceof applied stress, the effect of strain rate, and the developmentof thermo-induced nonequilibrium grain boundary segregationtheory.

ACKNOWLEDGMENTThe authors gratefully appreciate the financial supports from

the Alexander von Humboldt Foundation (AvH), the NationalNatural Science Foundation of China (NSFC) (No. 51001011),the Deutscher Akademischer Austausch Dienst (DAAD), theDeutsche Forschungsgemeinschaft (DFG) (No. SCHM 1182/9),the Research Fund for the Doctoral Program of Higher Educa-tion of China (No. 20090006120011), the Beijing MunicipalNatural Science Foundation (No. 2123064), and the Fundamen-tal Research Funds for the Central Universities (No. FRF-TP-12-042A). The authors appreciate the enlightening discussionswith Prof. Tingdong Xu (Central Iron and Steel Research In-stitute, China), Prof. Dr. Pavel Lejcek (Academy of Sciences

of the Czech Republic), Prof. Dr. Ulrich Krupp (Universityof Applied Sciences Osnabruck, Germany) and Dr. DebashisMukherji (Technical University Braunschweig, Germany), andwould like to thank Prof. E. Nembach, Dr. D. Baither, and Dr.Z. Balogh (University of Muenster) for valuable suggestions.

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Documents

Mechanism of Intermediate Temperature Embrittlement of Ni and Ni-based Superalloys