Low temperature neutron irradiation effects on microstructure and tensile properties of molybdenum

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Journal of Nuclear Materials 376 (2008) 11–28

Low temperature neutron irradiation effects on microstructureand tensile properties of molybdenum

Meimei Li a,, M. Eldrup b, T.S. Byun a, N. Hashimoto c, L.L. Snead a, S.J. Zinkle a

a Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USAb Materials Research Department, Risø National Laboratory, Technical University of Denmark, DK-4000, Roskilde, Denmark

c Materials Science Division, Hokkaido University, Sapporo 060-8628, Japan

Received 5 June 2007; accepted 12 December 2007

Abstract

Polycrystalline molybdenum was irradiated in the hydraulic tube facility at the High Flux Isotope Reactor to doses ranging from7.2 � 10�5 to 0.28 dpa at �80 �C. As-irradiated microstructure was characterized by room-temperature electrical resistivity measure-ments, transmission electron microscopy (TEM) and positron annihilation spectroscopy (PAS). Tensile tests were carried out between�50 and 100 �C over the strain rate range 1 � 10�5 to 1 � 10�2 s�1. Fractography was performed by scanning electron microscopy(SEM), and the deformation microstructure was examined by TEM after tensile testing. Irradiation-induced defects became visible byTEM at �0.001 dpa. Both their density and mean size increased with increasing dose. Submicroscopic three-dimensional cavities weredetected by PAS even at �0.0001 dpa. The cavity density increased with increasing dose, while their mean size and size distribution wasrelatively insensitive to neutron dose. It is suggested that the formation of visible dislocation loops was predominantly a nucleation andgrowth process, while in-cascade vacancy clustering may be significant in Mo. Neutron irradiation reduced the temperature and strainrate dependence of the yield stress, leading to radiation softening in Mo at lower doses. Irradiation had practically no influence on themagnitude and the temperature and strain rate dependence of the plastic instability stress.� 2008 Published by Elsevier B.V.

1. Introduction

Molybdenum is of great interest for high temperatureapplications in advanced fission and fusion reactor systemsbecause of its high melting point, excellent high tempera-ture strength, good thermal conductivity, resistance to irra-diation-induced swelling and corrosion resistance in liquidmetal coolants [1]. However, Mo, like other body-centeredcubic (bcc) metals, is susceptible to low temperaturesembrittlement and suffers an increase in ductile-brittle tran-sition temperature (DBTT) after neutron exposure [2–11].The improvement in low temperature ductility of neu-tron-irradiated Mo is of great importance for its applica-tions in advanced nuclear systems.

0022-3115/$ - see front matter � 2008 Published by Elsevier B.V.

doi:10.1016/j.jnucmat.2007.12.001

Corresponding author. Present address: Argonne National Laboratory.Tel.: + 1 630 2525111; fax: + 1 630 2523604.

E-mail address: [email protected] (M. Li).

Low temperature embrittlement is associated with irradi-ation hardening that is controlled by the interactions ofmobile dislocations and irradiation-induced defects.Whether or not Mo can be engineered to resist irradiation-induced low temperature embrittlement is dependent onthe formation process of sessile defect clusters in Mo. It isknown that two formation processes are often competingduring irradiation, namely in-cascade clustering and diffu-sive nucleation and growth. If the formation of sessile defectclusters is dominated by a nucleation and growth processsuch as in bcc Fe [12–16], solute additions may have a strongeffect on the formation of defect clusters and radiationhardening. The resistance to irradiation embrittlementmay, therefore, be improved by metallurgical approaches.On the other hand, if large, sessile defect clusters originatefrom displacement cascades, e.g. in bcc W [12–14,17], thelow temperature irradiation embrittlement is essentially aninherent problem, and may not be overcome.

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12 M. Li et al. / Journal of Nuclear Materials 376 (2008) 11–28

Defect formation and accumulation in irradiated mate-rials is influenced largely by crystalline structure andatomic mass. Molecular dynamic (MD) simulations andheavy-ion irradiation experiments have clearly showed thatdisplacement cascades are more compact and productionefficiency of visible defect clusters is greater in fcc Cu thanin bcc Fe [12,13,18–21]. The difference is primarily due tothe more compact lattice structure of fcc metals vs. themore open lattice structure of bcc metals. The mass differ-ence between metals is important in the defect productionand accumulation as well. The displacement cascades werefound to be more compact, and the energy threshold forsubcascade formation is higher in a high-mass metal thanin a medium-mass metals, such as bcc Mo vs. bcc Fe[20,21]. These differences in displacement cascades have sig-nificant implications in in-cascade clustering and defectcluster evolution.

While the fundamentals of defect production and evolu-tion in medium-mass metals, such as fcc Cu, Ni and bcc Feare well established, the understanding of defect produc-tion in high-mass metals such as fcc Pd and bcc Mo andW is still limited. Recent MD computer simulations ofatomic displacement cascades in Mo [20] have providednew insights into the primary radiation damage in Mo,and have identified the differences in in-cascade clusteringbehavior between medium-mass bcc Fe and high-massbcc Mo. Validation of these simulation results requiresexperimental evidence from well-designed mechanistic neu-tron irradiation experiments such as low temperature, lowdose neutron irradiation. One of the primary interests ofthis study is to provide direct measurements of defect clus-ters and resultant hardening to validate in-cascade cluster-ing results obtained from MD simulations.

The interaction between irradiation-induced defect clus-ters and glide dislocations is the key to understand thehardening and embrittlement in irradiated Mo. Irradiationhardening not only depends on the nature and size ofdefect clusters effective as barriers to dislocation motions,but also can be strongly affected by the strain rate and testtemperature [22–30]. In contrast to fcc crystals, a large tem-perature and strain rate dependence of yielding occurs inbcc metals even in unirradiated conditions. The yield stressincreases rapidly with decreasing temperature and increas-ing strain rate in unirradiated Mo [31]. A good understand-ing of the radiation hardening mechanism in bcc Morequires detailed studies of neutron irradiation effects onthe temperature and strain rate dependence of the flowstress. Temperature and strain rate-insensitive irradiationhardening implies that irradiation does not affect therate-controlling deformation mechanism operative in theunirradiated condition, i.e. athermal hardening resultedfrom long-range barriers, while temperature and strainrate-sensitive irradiation hardening indicates that irradia-tion-induced obstacles can be overcome by dislocationswith thermal assistance, i.e. thermal hardening fromshort-range barriers [23]. The study on the dependence ofthe true stress at the ultimate tensile strength, so-called

plastic instability stress (PIS) on the temperature and thestrain rate is lacking.

The increase in yield stress is often accompanied by aloss of strain hardening capacity and tensile ductility, andcan also lead to embrittlement by shifting the DBTT to ahigher temperature. Plastic deformation tends to occur inan inhomogeneous manner and to be localized in disloca-tion channels after irradiation. Premature plastic instabilityoccurs at yield once the yield strength reaches the level ofthe PIS of an irradiated material [32–35]. Wechsler [36] sug-gested that the low temperature irradiation embrittlementis likely associated with changes in plastic properties, par-ticularly with inhomogeneous plastic deformation ratherthan changes in inherent fracture processes. Crack forma-tion from coarse slip steps and the intersection of coarseslip bands was suggested to cause ‘channel fracture’[37,38]. Cracks may also initiate at grain boundaries dueto local stress concentration caused by dislocation pile-upin a defect-free channel [5]. Direct correlation between dis-location channeling and premature plastic instability andcorrelation between dislocation channeling and radiationembrittlement, however, have not been confirmed byexperiments.

The experiments described in this paper were designedto understand the role of in-cascade clustering in the for-mation of sessile defect clusters and the nature of interac-tions between irradiation-induced defects and movingdislocations. The ultimate purpose is to determine whetheror not the resistance to low temperature embrittlement ofirradiated Mo can be improved by metallurgicalapproaches. Neutron irradiation experiments were per-formed at seven low doses over a range of 7.2 � 10�5 to0.28 dpa at reactor ambient temperature (�80 �C) on pureMo. The formation of sessile defect clusters was investi-gated by characterizing as-irradiated microstructure byelectrical resistivity, transmission electron microscopy(TEM), and positron annihilation spectroscopy (PAS).Irradiation hardening mechanisms in Mo were understoodby examining the dose, temperature and strain rate depen-dence of the yield stress from tensile properties measure-ments between �50 and 100 �C at strain rates over therange of 1 � 10�5 to 1 � 10�2 s�1. Deformation micro-structure of irradiated specimens was examined after tensiletesting. Fracture surfaces of broken tensile specimens wereobserved by scanning electron microscopy (SEM). Theeffects of neutron irradiation on strain hardening, plasticinstability and fracture in Mo was described, and deforma-tion and fracture mechanisms in neutron-irradiated Mowere discussed.

2. Experimental procedure

The material examined was low-carbon arc-cast(LCAC) Mo with purity >99.95% and interstitial impuritycontents of 140 wppm C, 6 wppm O and 8 wppm N. Sub-size SS-3 sheet tensile specimens with gauge dimensionsof 7.62 � 1.52 � 0.50 mm and rectangular coupon speci-

Table 1Tensile test conditions

Dose (dpa) TestTemperature(�C)

Strain rate(s�1)

Dose(dpa)

TestTemperature(�C)

Strainrate (s�1)

Unirr. �100 1 � 10�3 Unirr. 22 1 � 10�1

�50 1 � 10�3 22 1 � 10�2

�25 1 � 10�3 22 1 � 10�3

22 1 � 10�3 22 1 � 10�4

100 1 � 10�3 22 1 � 10�5

7.2 � 10�5 �50 1 � 10�3 0.003 22 1 � 10�2

�25 1 � 10�3 22 1 � 10�3

22 1 � 10�3 22 1 � 10�4

100 1 � 10�3 22 1 � 10�5

7.2 � 10�4 �50 1 � 10�3 0.029 22 1 � 10�3

22 1 � 10�3 22 1 � 10�4

100 1 � 10�3 22 1 � 10�5

0.0072 �50 1 � 10�3

�25 1 � 10�3

22 1 � 10�3

100 1 � 10�3

0.072 �50 1 � 10�3

�25 1 � 10�3

22 1 � 10�3

100 1 � 10�3

0.28 22 1 � 10�3

100 1 � 10�3

10-5

10-4

10-3

10-2

10-1

100

50

60

70

80

90

Mo Neutron-irradiated at 80oC

ρ (n

Ω-m

)

Dose (dpa)

Unirradiated

Fig. 1. Room-temperature electrical resistivity of neutron-irradiated Moas a function of neutron dose. The mean value of electrical resistivity forunirradiated Mo is indicated by the dash line.

M. Li et al. / Journal of Nuclear Materials 376 (2008) 11–28 13

mens with dimensions of 25.4 � 4.95 � 0.25 mm weremachined from a cold-rolled 0.5 mm thick sheet. Couponspecimens were used for TEM microstructural character-ization and PAS measurements. Following machining,specimens were annealed at 1200 �C for 1 h in high vac-uum, resulting in an equiaxed grain structure with a grainsize of 70 lm.

Irradiation of specimens was carried out in the hydraulictube facility in the High Flux Isotope Reactor at the OakRidge National Laboratory. Specimens were loaded in per-forated rabbit capsules allowing contact with the flowingcoolant in the hydraulic tube to maintain the specimentemperature at �80 �C. Specimens were irradiated to neu-tron fluences in the range of 2 � 1021 to 8 � 1024 n/m2

(E > 0.1 MeV), corresponding to displacement damagelevels 7.2 � 10�5, 7.2 � 10�4, 7.2 � 10�3, 0.072 and0.28 dpa. Unexpected weight loss and thickness reductionwas observed on the specimens irradiated to high neutronfluences due to corrosion of molybdenum in coolant water.Two additional perforated rabbit capsules were irradiatedto ensure water reaction had no effect on the results. In thisadditional irradiation experiment, tensile specimens wereindividually wrapped in thin high-purity Al foil and vac-uum-sealed by electron-beam welding to minimize thermalgradients between the specimens and the flowing water.The thin-foil-wrapped specimens were then irradiated at�80 �C to neutron fluences of 8.4 � 1022 n/m2 and8.1 � 1023 n/m2 (E > 0.1 MeV), corresponding to displace-ment damage doses of 0.003 and 0.029 dpa. Post-irradia-tion measurements of electrical resistivity, as-irradiatedmicrostructure and tensile properties showed consistentresults with data obtained in the first experiment.

As-irradiated microstructure was examined by TEM tocharacterize defect cluster density, mean size and size distri-bution. The 3 mm discs were punched from coupon speci-mens or the grip regions of tensile specimens. Discspecimens were electropolished in a Tenupol twin-jet pol-ishing unit using a solution of 7:1 methanol and sulphuricacid cooled to �10 �C. Microstructure examination wasperformed using combined bright field (BF) and weakbeam dark field (WBDF) imaging techniques. Weak beamdark field imaging conditions were at (g/5g or g/6g,g = 110) near a zone axis of ½�111�. The defect cluster den-sity and size measurements were made from WBDFimages. The defect density measurements were checkedby plotting the areal density vs. foil thickness, and the vol-umetric density was determined by the slope of the plot[15].

Positron annihilation spectroscopy measurements wereconducted on 5 � 5 mm specimens cut from coupon speci-mens for each neutron dose. The specimens were electropo-lished before the measurements to remove surface effects. Aconventional positron lifetime spectrometer was used in themeasurements, and the details of the technique are given inRefs. [39,40]. All specimens were measured several times toassure reproducibility. The results for each dose are basedon lifetime spectra that typically contain a total of 12–

16 � 106 counts. The recorded spectra were evaluated usingthe PALSfit program [41].

Electrical resistivity was measured on tensile specimensat room temperature before and after irradiation using afour-point probe technique following the guidelines inASTM Standard Test Method for Resistivity of ElectricalConductor Materials ASTM B 193-02.

Tensile tests were conducted on unirradiated and irradi-ated specimens between �50 and 100 �C in air or a mixed


air and cold nitrogen environment at a strain rate over arange from 1 � 10�5 to 1 � 10�1 s�1. The tensile test condi-tions are given in Table 1. Load and displacement datawere recorded and used to determine stress–strain curves

Fig. 2. Weak beam dark field images of defect clusters at different dose levels taof dislocation loops in a row) were observed at higher doses.

and tensile properties. Fractographic examinations wereperformed by SEM after tensile testing. The SEM imagestaken at low magnification (30�–50�) were used to deter-mine final areas of the cross-section, and the reduction in

ken with a ½�111� zone axis and g = [110] (g/5g, 6g). Rafts (aligned groups

100

1000

Tirr= 50-100oC

022

/m3 )

LCAC Mo, this studyTanaka Mo Singh Mono Mo Singh Mono MoRe Singh Poly MoRe Singh TZM Muller Mo, proton Muller TZM, proton Muller Mo-5Re, proton


area (RA) was calculated as the percent change in cross-sectional area at failure.

The deformed microstructure of irradiated specimenswas examined by TEM in the uniformly strained gaugesection of broken tensile specimens.

10-4

10-3

10-2

10-1

100

0.1

1

10

Clu

ster

den

sity

(x1

Dose (dpa)

Niebel, Mo

5

6

7

8

9

10

Tirr= 50-100oC

ean

siz

e (n

m)

LCAC Mo, this study Tanaka Mo Singh Mono Mo Singh Mono MoRe Singh Poly MoRe Singh TZM Muller Mo, proton Muller TZM, proton Muller Mo-5Re, proton Niebel, Mo

3. Results

3.1. As-irradiated microstructure

3.1.1. Room-temperature electrical resistivity

The dose dependence of room-temperature electricalresistivity for neutron-irradiated Mo is shown in Fig. 1.The data points for each dose are the mean values of fivemeasurements from four unique specimens. The meanvalue of electrical resistivity for unirradiated Mo is indi-cated by the dash line. No measurable change in electricalresistivity was observed after irradiation to the lowest dose,7.2 � 10�5 dpa. An initial low rate of resistivity increasewas followed by a high rate between �0.01 and 0.1 dpa,followed by an asymptotic approach to a saturation valueat doses above 0.1 dpa. Calculations showed that the influ-ence of the neutron-induced transmutation products onresistivity change in the irradiated specimens was insignifi-cant. Detailed discussion is given elsewhere [42].

10-5

10-4

10-3

10-2

10-1

100

0

1

2

3

4

Clu

ster

m

Dose (dpa)

Fig. 3. Dose dependence of (a) defect cluster density and (b) defect clustermean size in neutron-irradiated Mo (solid lines show the data of thisstudy).

3.1.2. Microstructure characterization by TEM

Fig. 2 shows representative WBDF images for Mo irra-diated in the dose range from 7.2 � 10�5 to 0.28 dpa. Thedefect cluster density and mean size are plotted as a func-tion of dose in Fig. 3, and the size distribution are plottedas the histograms with 0.5 nm intervals in Fig. 4. Fig. 3 alsoincludes published data on pure Mo and Mo alloys thatwere neutron-irradiated at 50–100 �C for comparison [9–11,43,44]. As shown in Fig. 2, the as-irradiated microstruc-ture consisted of uniformly-distributed dislocation loops atlower doses. The nature of the loops could not be resolveddue to their small sizes. Previous studies on Mo annealedfollowing irradiation have found that the dislocation loopsare predominantly interstitial-type and have the Burgersvectors b = a/2h111i [45–50]. A highly-aggregated loopstructure, so-called rafts, was observed at higher doses.Raft formation was detectable in the 0.0072 dpa specimen,indicated by aligned groups of two or three dislocationloops in a row. The formation of rafts became more pro-nounced as dose increased, and the raft length increasedsignificantly with increasing dose.

A strong dose dependence of defect cluster density wasobserved in Mo. No visible defect clusters were observedin the 7.2 � 10�5 dpa specimen. The visible cluster densityinitially increased with increasing dose, reached a maxi-mum at 0.029 dpa, and then decreased with increasingdose. The mean cluster size was dependent on irradiationdose as well. The mean cluster size increased linearly withdose from 1.94 to 3.36 nm between 7.2 � 10�4 and0.28 dpa. The cluster size distribution showed two charac-

teristic features: a shift of peaks to the large cluster size anda longer tail at large cluster size with increasing doses.

Weak contrast resembling cavities was revealed at highmagnifications (300k�) only in the 0.28 dpa specimen.Detailed characterization of cavities was difficult due toweak contrast. The estimated cavity density from TEMmicrographs was on the order of 1023 m�3 with a mean sizenear the TEM resolution limit (�1 nm).

3.1.3. Microstructure characterization by PAS

Positron annihilation spectroscopy is particularly sensi-tive to the size and density of vacancy type defects, typi-cally in the size range from mono-vacancies up to nano-voids containing �50–100 vacancies and of densities ofthe order of parts per million. As discussed in some detailin the reference [51], it is possible to estimate size distribu-tions of nanometer-sized cavities in irradiated materials.Based on such estimates, the total cavity density and the

0 1 2 3 4 5 6 7 8 9 10

00

1 2 3 4 5 6 7 8 9 10 0 1 2 3 4 5 6 7 8 9 10

0 1 2 3 4 5 6 7 8 9 100

10

20

30

40

50

0

0

10

20

30

40

50

Def

ect

clu

ster

fra

ctio

n (

%)

Def

ect

clu

ster

fra

ctio

n (

%)

Def

ect

clu

ster

fra

ctio

n (

%)

Def

ect

clu

ster

fra

ctio

n (

%)

Defect cluster size (nm)


Defect cluster size (nm) Defect cluster size (nm)



0.00072 dpa

mean size = 1.94 nmmean size = 2.24 nm

0.003 dpa

10

20

30

40

50

mean size = 2.36 nm

0.0072 dpa

10

20

30

40

50

mean size = 2.71 nm

0.029 dpa

Def

ect

clu

ster

fra

ctio

n (

%)

0 1 2 3 4 5 6 7 8 9 100

10

20

30

40

50

mean size = 2.95 nm

0.072 dpa

0 1 2 3 4 5 6 7 8 9 100

10

20

30

40

50

mean size = 3.36 nm

0.28 dpa

Def

ect

clu

ster

fra

ctio

n (

%)

Fig. 4. The size distribution of TEM-visible defect clusters in neutron-irradiated Mo.


mean diameter can be calculated as a function of dose. Thepositron lifetime spectra for Mo were analyzed on theassumption that the distribution of lifetimes can be approx-imated by four bins: 190, 260, 350 and 470 ps, equivalent tothree-dimensional vacancy clusters containing 1, 3, 10 and50 or more vacancies, respectively. The calculated size dis-tributions for the vacancy clusters are shown in Fig. 5(a),and the total cavity density and the mean diameter as afunction of dose are given in Figs. 5(b) and (c). Thevacancy clusters were present in Mo irradiated at all dosesfrom 7.2 � 10�5 to 0.28 dpa. Their size distribution wasnearly unchanged during irradiation. The overall densityof vacancy defects increased with dose by more than a fac-

tor of ten in the examined dose range with a tendency ofleveling off at the highest dose. When the number densityof cavities containing greater than 50 vacancies was plottedseparately, the density increased from �1021 m�3 to�1023 m�3 as dose increased from 0.0072 dpa to 0.28 dpa.The number density of these large cavities in the 0.28 dpaspecimen is comparable to the estimate from TEM obser-vations. No such large cavities were observed in the0.00072 dpa specimen, and the data point for the7.2 � 10�5 dpa specimen cannot be explained. It is alsonoted that the cavity density follows a linear relation withneutron dose for 50-plus vacancy clusters and an exponentof �0.5 for smaller vacancy clusters. The mean size of

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.70.0

0.2

0.4

0.6

0.8

1.0

0.000072 dpa 0.00072 dpa 0.0072 dpa 0.072 dpa 0.28 dpa

Vac

ancy

clu

ster

fra

ctio

n

Size (nm)

10-5 10-4

10-3

10-2

10-1

100

10-4

10-3

10-2

10-1

100

101

102

total density mono-vacancy 3 vac. 10 vac. >=50 vac.

Nu

mb

er d

ensi

ty (1

024

/m3

)

Dose (dpa)

n = 1

n = 0.5

Nv = A(dose)n

10-5

10-4

10-3

10-2

10-1

100

0.0

0.2

0.4

0.6

0.8

1.0

Mea

n D

iam

eter

(nm

)

Dose (dpa)

vacancy clusters

mono-vacancies and vacancy clusters

Fig. 5. (a) Size distribution, (b) dose dependence of number density, and(c) dose dependence of mean sizes of mono-vacancies and vacancy clustersin Mo neutron-irradiated at 80 �C.


vacancy clusters were calculated with and without includ-ing mono-vacancies, and the mean size showed only a smallincrease with dose, if any.

3.2. Post-irradiation tensile properties

3.2.1. Stress–strain behavior

The engineering stress–strain curves are shown in Fig. 6for unirradiated and irradiated Mo tested over the temper-ature range of �50 to 100 �C at a nominal strain rate of1 � 10�3 s�1. Most of these results were reported previously[52]. The true stress–strain curves were calculated from theengineering stress–strain curves in the uniform elongationrange, and they are given in Fig. 7. The yield point behavior,i.e. the upper and lower yield points and Luders strain, wasobserved in the unirradiated specimens and specimens irra-diated to low and intermediate doses (e.g. <0.01 dpa). Thepresence of the yield point and the magnitude of yield dropwere dependent on both test temperature and neutron dose.The yield point was less evident as dose increased, and morepronounced as the test temperature decreased. The low andintermediate dose specimens also exhibited extensive strainhardening and uniform deformation at all test tempera-tures. The true uniform elongation was still about 10% atthe test temperature of �50 �C. The plastic instability stress(PIS), i.e. the true stress at the ultimate tensile strength, wasapparently independent of irradiation dose at a given testtemperature (see Fig. 7). The strain hardening capacity ofneutron-irradiated Mo was completely lost at doses>�0.01 dpa, leading to the onset of plastic instability atyield and a nearly zero uniform elongation.

Fig. 8 shows the engineering stress–strain curves forunirradiated Mo and Mo irradiated to 0.003 and0.029 dpa tested over the strain rate range of 1 � 10�5 to1 � 10�2 s�1 at room temperature. The sensitivity of theyield point behavior to the strain rate is well-illustrated inFig. 8(a) for unirradiated Mo and in Fig. 8(b) for Mo irra-diated to 0.003 dpa. The upper and lower yield point andLuders elongation were more pronounced as the strain rateincreased. Neutron irradiation reduced the yield pointeffect. Increasing strain rates increased the tensile strengthand decreased the tensile ductility. Molybdenum irradiatedto 0.029 dpa (Fig. 8(c)) showed completely brittle behavioreven at a strain rate of 1 � 10�5 s�1.

3.2.2. Tensile properties

Tensile properties, i.e. yield stress (YS), plastic instabil-ity stress (PIS), uniform elongation (UE), total elongation(TE) and reduction in area (RA) are shown in Fig. 9 inthree-dimensional plots for irradiated Mo tested at a strainrate 1 � 10�3 s�1 between �50 and 100 �C to illustrate thedose and test temperature dependence of tensile properties.The yield stress was taken as the lower yield point whensignificant strain hardening and homogeneous deformationwere observed; the yield stress was taken as the upper yieldpoint when the onset of plastic instability at yield occurred.

As shown in Fig. 9(a), Mo experienced radiation-induced softening (decrease in yield stress) at lower doseswhen tested at low temperature; only radiation hardening(increase in yield stress) occurred at higher doses. The soft-ening effect was stronger at lower test temperatures. No

0 10 20 30 40 50 60 700

200

400

600

800

1000

1200Tirr=80oC, Ttest=100oC

Str

ess

(MP

a)

0

200

400

600

800

1000

1200

Str

ess

(MP

a)

0

200

400

600

800

1000

1200

Str

ess

(MP

a)S

tres

s (M

Pa)

Strain (%) Strain (%)0 10 20 30 40 50 60 70

0.28

dpa

0.07

2 dp

a

7.2x

10-3

dpa

7.2x

10-4

dpa

Uni

rrad

iate

d

7.2x

10-5

dpa

Tirr=80oC, Ttest=22oC

0.02

9 dp

a

7.2x

10-3 d

pa

0.28

dpa

0.07

2 dp

a

3.0x

10-3 d

pa

7.2x

10-4 d

pa

Uni

rrad

iate

d

7.2x

10-5 d

pa

0 10 20 30 40 50

Tirr=80oC, Ttest=-25oC

7.2x

10-3 d

pa

0.07

2 dp

a

Strain (%)0 10 20 30 40 50

Strain (%)

Uni

rrad

iate

d

7.2x

10-5 d

pa

0

200

400

600

800

1000

1200Tirr=80oC, Ttest=-50oC

7.2x

10-3 d

pa

0.07

2 dp

a

7.2x

10-4 d

pa

Uni

rrad

iate

d

7.2x

10-5 d

pa

Fig. 6. Engineering stress–strain curves for Mo neutron-irradiated at 80 �C and tested at: (a) 100 �C, (b) 22 �C, (c) �25 �C and (d) �50 �C at a strain rateof 1 � 10�3 s�1.


irradiation softening was observed by the tests at 100 �C.The softening effect was more evident in the 7.2 � 10�4 dpaspecimen than in the 7.2 � 10�5 dpa specimen at a giventest temperature, e.g. a 71 MPa drop in the yield stress at7.2 � 10�5 dpa vs. 147 MPa drop at 7.2 � 10�4 dpa at thetest temperature of �50 �C. When irradiation hardeningoccurred, the yield strength increased significantly withincreasing dose up to �0.1 dpa. Saturation or a decreasein the yield strength was observed above �0.1 dpa. Asdescribed in the previous section, the PIS was not depen-dent on irradiation dose, while the PIS increased signifi-cantly with decreasing test temperature (shown inFig. 9(b)). Note that the PIS is not defined when prematureplastic instability at yield occurred.

Irradiation hardening was accompanied by a loss inductility. As shown in Fig. 9(c), the uniform elongationwas reduced to nearly zero at 0.072 dpa at all test tem-peratures, due to either plastic instability at yield athigher test temperatures or embrittlement at lower testtemperatures. The critical dose for plastic instability atyield was �0.072 dpa. When premature plastic instabilityoccurred, the total elongation remained moderate (�5%)(see Fig. 9(d)) and the reduction in area was nearlyunchanged within data scatter (see Fig. 9(e)). Embrittle-ment was more severe at lower test temperatures, and

both the total elongation and reduction in area droppeddramatically.

The strain rate dependence of tensile properties is shownin Fig. 10 for Mo tested at room temperature. Increasingstrain rates led to increases of both the yield stress andthe PIS and decrease in ductility. The yield stress showeda stronger dependence on the strain rate than the PIS forunirradiated Mo. Irradiation to 0.003 dpa reduced thestrain rate dependence of the yield stress, but had noimpact on the magnitude and strain rate dependence ofthe PIS. The reduced strain rate dependence of the yieldstress resulted in softening at higher strain rates, and hard-ening at lower strain rates in the 0.003 dpa specimens. Thecrossover of softening to hardening occurred at3 � 10�3 s�1. The strain rate dependence of the yield stressand the PIS was almost the same at 0.003 dpa. Both theyield stress, rys and plastic instability stress, rpis exhibiteda linear dependence on the strain rate, _e in semi-log scale,and the fitting of the YS data and the PIS data gave that:

rys ðMPaÞ ¼ 711:3þ 82:3 log _e unirradiated;

rys ðMPaÞ ¼ 617:3þ 45:1 log _e 0:003 dpa;

rpis ðMPaÞ ¼ 775:5þ 43:9 log _e

for both unirradiated and 0:003 dpa:

ð1Þ

0.00 0.05 0.10 0.15 0.20 0.25 0.300

200

400

600

800

1000

1: Unirradiated2: 0.000072 dpa3: 0.00072 dpa4: 0.0072 dpa5: 0.072 dpa6: 0.28 dpa

Tirr=80oC, Ttest=100oCT

rue

stre

ss (

MP

a)T

rue

stre

ss (

MP

a)

0

200

400

600

800

1000

Tru

e st

ress

(M

Pa)

Tru

e st

ress

(M

Pa)

True strain True strain

321

4

5

6

PIS

0.00 0.05 0.10 0.15 0.20


4

32

568

7

1: Unirradiated2: 0.000072 dpa3: 0.00072 dpa4: 0.003 dpa5: 0.0072 dpa6: 0.029 dpa7: 0.072 dpa8: 0.28 dpa

1

PIS

0.00 0.05 0.10 0.15 0.200

200

400

600

800

1000


1: Unirradiated2: 0.000072 dpa3: 0.0072 dpa4: 0.072 dpa

4

True strain True strain

1 2 3 PIS

0.00 0.05 0.10 0.15 0.200

200

400

600

800

1000


1: Unirradiated2: 0.000072 dpa3: 0.00072 dpa4: 0.0072 dpa5: 0.072 dpa

12

3

4

5

PIS

Fig. 7. True stress-strain curves for Mo neutron-irradiated at 80 �C and tested at: (a) 100 �C, (b) 22 �C, (c) �25 �C and (d) �50 �C at a strain rate of1 � 10�3 s�1.


The 0.029 dpa specimens exhibited completely brittlebehavior at all strain rates.

3.3. Fractography of neutron-irradiated Mo

Fig. 11 shows SEM images of the fracture surfaces forMo specimens neutron-irradiated to 7.2 � 10�5 dpa andtested at 1 � 10�3 s�1 between �50 and 100 �C. Their cor-responding tensile curves are also shown in Fig. 11. Thefracture mode at this dose level depends strongly on testtemperature. A complete ductile, layered structure (delam-ination) was observed when the specimen was tested at100 �C. As test temperature decreased, the fraction of areasexhibiting ductile failure decreased, and the amount ofcleavage and grain boundary brittle fracture increased.At the test temperature of �50 �C, failure was solely bycleavage and (minor) grain boundary fracture.

A set of SEM fractographs and their corresponding ten-sile curves are shown in Fig. 12 for the specimens neutron-irradiated to 0.072 dpa. The dose dependence of fracturemode can be revealed by comparing Figs. 11 and 12. Asignificant difference in fracture mode was seen in thespecimens tested at 22 �C. The 7.2 � 10�5 dpa specimenshowed a dominant ductile fracture mode mixed with slight

cleavage and grain boundary fracture, while the fracturemode in the 0.072 dpa specimen was dominated by cleav-age fracture with a small fraction of layered, ductile frac-ture and grain boundary fracture.

3.4. Deformed microstructure of neutron-irradiated Mo

Representative TEM micrographs showing deformedmicrostructure in the uniform strain region of neutron-irra-diated Mo are presented in Fig. 13 along with the corre-sponding engineering stress–strain curves. No dislocationchannels were observed in the irradiated and tensile-deformed Mo specimens that exhibited extensive strainhardening. Only dislocation networks were observed in thiscase. Dislocation channels were observed in the 0.0072 dpaspecimen tested at �50 �C when it failed at yield (0.24%plastic elongation) (micrograph shown in Fig. 13(a) andtensile curve indicated by No. 1 in Fig. 13(c)). The micro-structure of this specimen consisted of nearly defect-freechannels and undeformed regions containing a high num-ber density of irradiation-induced defect clusters compara-ble to as-irradiated Mo. The channels were about 80 nmwide uniformly and the channel walls were smooththroughout the channel. Dislocation channels were also

0 10 20 30 40 50 600

200

400

600

800

1000

5

4

32 5: 1x10-5 s-1

4: 1x10-4 s-13: 1x10-3 s-12: 1x10-2 s-1

Unirradiated Mo, Ttest=22oC

Str

ess

(MP

a)

0

200

400

600

800

1000

Str

ess

(MP

a)

0

200

400

600

800

1000

Str

ess

(MP

a)

Strain (%)

0 10 20 30 40 50 60

Strain (%)

1: 1x10-1 s-1

1

4

3

2

1

Tirr=80oC 0.003 dpa, Ttest=22oC

1: 1x10-2 s-1

2: 1x10-3 s-1

3: 1x10-4 s-1

4: 1x10-5 s-1

0 2 4 6 8 10 12 14 16 18 20

Strain (%)

1x10-3 s-1

1x10-5 s-1

1x10-4 s-1

Tirr=80oC 0.029 dpa, Ttest=22oC

Fig. 8. Engineering stress–strain curves at various strain rates for Mo (a)unirradiated, (b) irradiated to 0.003 dpa, and (c) irradiated to 0.029 dpatested at room temperature.


observed in the 0.072 dpa specimen tested at 22 �C (micro-graph shown in Fig. 13(b) and tensile curve indicated byNo. 2 in Fig. 13(c)). Channels were cleared of defects with

a few residual dislocations and debris from partially anni-hilated defect clusters. The width of channels varied (60–160 nm). Compared to the as-irradiated microstructure,the deformed matrix had a high density of dislocation linesdecorated with dislocation loops. Rafts seen in the as-irra-diated specimens were still present in irradiated anddeformed Mo. In general, thinner channels were observedat a lower test temperature. Due to the limited thin areain these specimens, complete quantitative characterizationof channels could not be performed.

4. Discussion

4.1. Formation of radiation-induced defect clusters

Microstructural evolution in Mo neutron-irradiated atlow doses at low temperature has been characterized byTEM, PAS and room-temperature electrical resistivity.Both dislocation loops and cavities were detected. TheTEM-visible dislocation loops are assumed to be predom-inantly interstitial-type according to previous loop analysesof neutron-irradiated Mo in the literature [45–50]. Severalsalient characteristics of defect clusters are summarized asfollows: (1) defect clusters were not visible by TEM at avery low dose of �0.0001 dpa; (2) the number density ofvisible defect clusters increased with increasing dose in asublinear manner, reached a maximum at �0.03 dpa anddecreased with continuous increase of dose, (3) the meansize of visible defect clusters increased from 1.94 to3.36 nm as dose increased, and the size distribution wasbroader at higher doses; (4) there was a strong indicationof aggregation of dislocation loops and formation of raftsat higher doses; (5) cavities were detected by PAS at alldoses between 7.2 � 10�5 and 0.28 dpa; (6) the mean sizeof cavities increased only slightly (if any) with increasingdoses; (7) weak contrast of cavities were observed byTEM only in the highest dose (0.28 dpa) specimen; (8)the estimated cavity density (�1023 m�3) by TEM is com-parable to the density of cavities containing more than 50vacancies detected by PAS; (9) there was no measurableincrease in room-temperature electrical resistivity at a verylow dose of �0.0001 dpa, and the electrical resistivityincreased continuously with increasing dose up to 0.28 dpa.

The microstructural features of dislocation loops andrafts are consistent with those observed in previous studiesof neutron-irradiated Mo at comparable irradiation tem-peratures. The defect cluster density was relatively low inthe present experiments compared to the results reportedby Tanaka et al. [9] and by Niebel and Wilkens [43] onpolycrystalline Mo (see Fig. 3). All three studies showeda continuous increase in number density below 0.03 dpa,while the decrease in density and size with further irradia-tion was more gradual in this study than that observed byTanaka et al. The accumulation rate of visible defect clus-ters reported by Niebel et al. was noticeably smaller thanall other studies due to the very high cluster density atlow doses. The material purity could be a major factor that

10-510

-410-310

-210-110

0 0

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1000

1200

-50-2522

YS

(MP

a)

Dose (dpa) 0

100

Temperature (o C)

10 -510 -4

10 -310 -2

10 -1

10 0

0200

400

600

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1000

1200

Temperature ( oC)

-50

-25

22

PIS

(M

Pa)

Dose (dpa)

100

0

10-5

10-4

10-3

10-2

10-1

100

0

10

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40

Temperature(o C)

-50

-25

22

Dose (dpa)

UE

(%)

0100

10-5

10-4

10-3

10-2

10-1

100

010

20

30

40

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60

-50

-25

22Temperature (oC)

TE

(%

)

Dos

e (d

pa)

0

100

0

20

40

60

80

100

10 -5

10 -4

10 -310 -2

10 -110 0

-50

-25

22

Temperature

( oC)

RA

(%

)

Dose (dpa)0

100

Fig. 9. Dose and temperature dependence of (a) yield stress (YS), (b) PIS, (c) uniform elongation (UE), (d) total elongation (TE) and (e) reduction in area(RA) for unirradiated and irradiated Mo tested at a strain rate 1 � 10�3 s�1 between �50 and 100 �C.


contributes to the variations in visible defect cluster densityamong different studies. The effect of impurity trapping ofself-interstitial atoms (SIA) in Mo has been demonstratedby Eyre et al. [47], where the visible clusters were signifi-cantly finer in low purity Mo.

The trends of cluster density and size changes with dosefor polycrystalline Mo in the present study are comparable

to the results on mono-crystalline Mo by Singh et al. [10].However, the cluster density in polycrystalline Mo wasnearly one order of magnitude higher, and the mean sizesof defect clusters were about a half that found in mono-crystalline Mo. Raft formation was observed at dosesabout 0.01 dpa and higher in both the mono-crystallineand polycrystalline Mo. Comparing pure Mo and Mo

10-5

10-4

10-3

10-2

10-1

200

300

400

500

600

700

800

900

1000

YS

/PIS

(M

Pa)

Strain rate (1/s)

YS, unirradiated PIS, unirradiatedYS, 0.003 dpa PIS, 0.003 dpaYS, 0.029 dpa

10-5

10-4

10-3

10-2

10-1

0

10

20

30

40

50

60

UE

or

TE

(%

)

Strain rate (1/s)

UE, unirradiated TE, unirradiated UE, 0.003 dpa TE, 0.003 dpa UE, 0.029 dpa TE, 0.029 dpa

Fig. 10. Strain rate dependence of (a) the yield stress (YS) and PIS and (b)uniform elongation (UE) and total elongation (TE) for unirradiated andirradiated Mo tested at room temperature. The 0.029 dpa specimensshowed premature brittle failure.


alloys, the cluster density in Mo–Re and TZM alloys washigher, and the cluster size increased more rapidly withincreasing dose [10,11,44]. Singh et al. [10,11] reported adecreased cluster density with increasing dose in the inter-mediate and high dose regime for neutron-irradiatedMo–Re and TZM alloys, while Muller [44] showed a nearlyconstant defect cluster density in Mo-Re and TZM alloysunder 600 MeV proton irradiation. It appears that theaddition of alloying elements affects the nucleation rateand mobility of defect clusters.

It is informative to compare the irradiated microstruc-ture between medium-mass bcc Fe and high-mass bcc Moneutron-irradiated near room temperature. The same typesof defect structure, i.e. dislocation loops, rafts and cavities,were observed in these two bcc metals [16,51]. Fig. 14 com-pares the dose dependence of loop and cavity densities andtheir mean sizes in Mo and Fe following neutron irradiationat 60–80 �C. The Fe data were taken from the Refs. [16,51].

No visible dislocation loops were observed in either Mo orFe at a very low dose (�0.0001 dpa) where displacementcascades were largely separated. The threshold dose for vis-ible loops was similar in both metals, �0.0007–0.0008 dpa.However, the loop density in neutron-irradiated Mo wasabout one order of magnitude higher than in Fe at0.001 dpa, �2 � 1022 m�3 in Mo vs. �1 � 1021 m�3 in Fe.The accumulation rate of visible loops in Mo and Fe wassimilar up to about 0.03 dpa where the loop density inMo reached a maximum and yet the loop density in Feincreased continuously with increasing dose up to �1 dpa.The cavity density is also slightly higher in Mo than inFe. The mean size of loops in Mo were larger than in Feat low doses but the loop size in Fe increased much morerapidly with dose than in Mo. The cavity size in Fe alsoshowed an evident increase, but not in irradiated Mo.

The lack of TEM-visible dislocation loops at a very lowdose indicated that large, sessile self-interstitial atoms(SIA) clusters were not produced directly in displacementcascades. Homogeneous nucleation and growth is insuffi-cient to form visible defect clusters at this low dose dueto its low rate and widely separated cascades. The sublineardefect accumulation behavior revealed by the dose depen-dence of visible loop density and the loop size dependenceupon dose at intermediate doses further indicate that thevisible interstitial loops were not formed by in-cascade clus-tering. Significant raft formation in Mo at higher dosesmay be explained by diffusional glide of small SIA clustersproduced in displacement cascades, which also explains thedecreased number density of visible defect clusters at higherdoses (>0.03 dpa). The reduced cluster density could be dueto destruction of pre-existing defects by cascade overlap-ping as well. Molecular dynamic simulations [20] havefound significant in-cascade interstitial clustering in irradi-ated Mo with a majority of interstitial defect clusters com-posed of only a few defects. These interstitial clusters takethe form of ½h111i primatic loop embryos and most areglissile at room temperature. Only a very low fraction ofclusters is in sessile configurations. A high number densityof large, sessile dislocation loops visible by TEM must beformed by both interactions between small glissile intersti-tial defects and through diffusive nucleation and growthprocesses during irradiation, similar to the cluster forma-tion process in irradiated Fe [16]. Despite the fact thatthe point defect nucleation and growth is dominant in theformation of large, sessile interstitial clusters in both Moand Fe, the higher number density and larger mean sizeof visible defect clusters at low doses in Mo imply thatin-cascade interstitial clustering in Mo may be more signif-icant than in Fe.

The findings of cavities in neutron-irradiated Mo at sucha low irradiation temperature, 80 �C (0.12 Tm, where Tm isthe melting point) in the present study are of great signifi-cance. It is noted that the irradiation temperature is consid-erably lower than the Mo recovery stage III temperature,150–200 �C (0.15–0.16 Tm) when vacancy migration occurs[53,54]. Positron annihilation spectroscopy detected cavi-

0 10 20 30 40 50 60 700

200

400

600

800

1000

1200

-50oC

-25oC

22oC

Str

ess

(MP

a)

Strain (%)

100oC

Tirr=80oC, 7.2x10-5 dpa

dε/dt = 1x10-3 s-1

Fig. 11. Fractography for Mo neutron-irradiated to 7.2 � 10�5 dpa and tested at 1 � 10�3 s�1 between �50 and 100 �C.

0 2 4 6 8 100

200

400

600

800

1000

1200

Tirr=80oC, 0.072 dpa

-25oC

-50oC

22oC

Str

ess

(MP

a)

Strain (%)

100oC

dε/dt = 1x10-3 s-1

Fig. 12. Fractography for Mo neutron-irradiated to 0.072 dpa and tested at 1 � 10�3 s�1 between �50 and 100 �C.


ties in Mo at a dose as low as �0.0001 dpa, similar to thethreshold dose in neutron-irradiated Fe [51]. The cavitydensities in Mo is slightly higher than in Fe, and the aver-age size of cavities in Mo is relatively insensitive to the neu-tron dose as opposed to an evident increase in cavity meansize with dose in Fe, as shown in Fig. 14 [16,51]. What ismore important is that the size distribution of vacancy clus-ter in Mo remained nearly constant with the peak at thesize of mono-vacancy, while the size distribution tends toshift with the peak at the increasing size with increasing

dose in Fe [51]. These observations suggest that in-cascadevacancy clustering is more significant in Mo than in Fe,which is supported by MD computer simulations [20,21].The increased tendency of in-cascade vacancy clusteringin Mo was also supported by the findings of higher defectyield (ratio of the number of visible vacancy loops and thenumber of cascades generated per unit area) and cascadeefficiency (ratio of the number of vacancies retained in aloop of average area and the calculated number of vacan-cies generated in the cascade) in Mo as compared to Fe

0 3 40

200

400

600

800

1000

1200curve 1: 0.0072 dpa/-50oCcurve 2: 0.072 dpa/22oC

Str

ess

(MP

a)

Strain (%)

1 2

1 2 5 6 7 8

Fig. 13. Deformed microstructure and their corresponding stress–straincurves for neutron-irradiated Mo (a) micrograph for the 0.0072 dpa/�50 �C specimen, (b) micrograph for the 0.072 dpa/22 �C specimen, (c)stress–strain curves of the 0.0072 dpa/�50 �C specimen and the 0.072 dpa/22 �C specimen.

10-5

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10-3

10-2

10-1

100

10-5

10-4

10-3

10-2

10-1

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1021

1022

1023

1024

1025

Mo loop density Fe loop density Mo cavity density Fe cavity density

Def

ect

Clu

ster

Den

sity

( /m

3 )

Dose (dpa)

Dose (dpa)

0.1

1

10Mo loop mean size Fe loop mean sizeMo cavity mean size Fe cavity mean size

Def

ect

Clu

ster

Mea

n S

ize

(nm

)

Fig. 14. Comparison of dose dependence of loop and cavity density (a)and mean size (b) in Mo and Fe irradiation at 60–80 �C. The Fe data weretaken from Refs. [16,51].


in heavy-ion irradiation experiments [12,13]. The averagedefect yield and cascade efficiency was 0.12 and 0.36,respectively for Mo irradiated with 60 keV self-ions atroom temperature, and yet no visible damage seen in bcc

Fe when irradiated with 80 keV self-ions. Although amajority of vacancy clusters predicted by MD simulationshad the sizes below TEM visibility limit, vacancy clustersas large as 60 defects were produced in a 50 keV cascadeenergy, and these sizeable clusters were considerably largerin Mo as compared to bcc Fe [20,21]. The morphology ofvacancy clusters observed in the present experiment is con-sistent with the predictions by MD simulations, i.e. all thelargest vacancy clusters are compact, three-dimensionalobjects.

4.2. Radiation softening and hardening

A strong effect of neutron irradiation on the yield stresswas found in Mo. The dose dependence of the change inyield stress clearly indicated three dose regimes (seeFig. 15): (1) radiation-induced yield softening (DrYS < 0)at lower doses (<�0.003 dpa), (2) radiation-induced yield

10-5 10-4 10-3 10-2 10-1 100-200

0

200

400

600Reduced yieldhardening

Yield hardening

Δσ ys

(MP

a)

Dose (dpa)

Ttest

=100oC, 1x10-3/s

Ttest

=22oC, 1x10-3/s

Ttest

=-25oC, 1x10-3/s

Ttest

=-50oC, 1x10-3/s

Ttest

=22oC, 1x10-2/s

Ttest

=22oC, 1x10-4/s

Ttest

=22oC, 1x10-5/s

Yield softening

Fig. 15. The change in the yield stress of neutron-irradiated Mo testedbetween �50 and 100 �C.


hardening (DrYS > 0) at intermediate doses (�0.003–0.1 dpa), (3) reduced radiation hardening (DrYS > 0 anddDrYS/d(dose) < 0) at higher doses (�>0.1 dpa). The cross-over of softening to hardening was dependent on the testtemperature and the strain rate. The softening effect wasstronger at a lower test temperature and a higher strainrate. The relation DrYS = A(/t)1/2 where (/t) is the neutronfluence, and A is a constant, holds up to �0.1 dpa if DrYS istaken as the increase in athermal stress of the yield stressafter irradiation.

Neutron irradiation had a pronounced effect on the tem-perature and strain rate dependence of the yield stress inMo as well (strain rate dependence shown in Fig. 10(a)and temperature dependence shown in Fig. 16). Asopposed to some irradiated fcc metals where the tempera-ture dependence of the yield stress is markedly increased

-80 -60 -40 -20 0 20 40 60 80 100 1200

200

400

600

800

1000

1200

1400 unirr 7.2x10-5 dpa 0.00072 dpa 0.0072 dpa

0.072 dpa

Temperature (oC)

YS

/PIS

(M

Pa)

Solid Symbols: YSOpen Symbols: PIS

Fig. 16. Temperature dependence of the yield stress (YS) and the plasticinstability stress (PIS) for unirradiated and irradiated Mo tested between�50 and 100 �C.

after neutron irradiation [22,55], the temperature andstrain rate dependence of the yield stress in Mo wasreduced by neutron irradiation. The temperature and strainrate dependence of the yield stress decreased as doseincreased up to �0.003 dpa, above which the temperatureand strain rate dependence of the yield stress remainednearly constant during further irradiation. It is interestingto note the temperature dependence of the yield stresswas stronger than that of the PIS in unirradiated Mo. Neu-tron irradiation had no effect on the magnitude and tem-perature dependence of the PIS, while irradiation reducedthe temperature dependence of the yield stress to nearlythe same level as the PIS before athermal hardeningoccurred.

Significant radiation softening and reduced test temper-ature and strain rate dependence of the yield stress at lowerdoses implies that radiation-induced defect clusters aremainly short-range thermal barriers to dislocation motion,affecting the rate-controlling deformation mechanism atlow temperatures. Previous analysis [52] showed that irra-diation-induced defects in Mo that account for short-rangebarriers are in agreement with the Fleischer theory andthey were mainly submicroscopic defects. A possible expla-nation for the reduced temperature dependence of the yieldstress could be due to trapping of interstitial solutes inpoint defect-impurity complexes.

The absence of a change in the temperature dependenceof the yield stress upon further irradiation (>�0.003 dpa)implies that irradiation-produced obstacles to dislocationmotion are largely athermal in nature [23]. A high numberdensity of visible dislocation loops may account for theselong-range barriers. This again suggests that large, sessiledefect clusters (athermal hardening centers) formed primar-ily through reactions between randomly diffusing mobiledefects. The hardening effect from large, sessile defect clus-ters may be understood using dispersed barrier hardeningmodels [56]. The barrier strength of defect clusters can bedetermined by correlating the athermal stress and the mea-sured density and size of defect clusters [42,52]. The calcu-lated value of barrier strength, a for visible defect clustersand large size cavities (containing 50 vacancies) increasedfrom 0.1 to 0.3–0.4 as doses increased from 0.003 to 0.3 dpa.

4.3. Plastic instability and fracture

As observed in several irradiated bcc metals [32], plasticinstability at yield (PIS = YS) occurs in Mo after irradia-tion to about 0.01–0.1 dpa, depending on the test tempera-ture and the strain rate. Below this critical dose for plasticinstability, the engineering stress–strain curves consist of afew distinct regions including elastic strain, yield drop,Luders strain, strain hardening, and necking; above thecritical dose, the curves show prompt ductile necking orcomplete brittleness. The prompt plastic instability hasbeen associated with the increase in yield stress and the lossof strain hardening capacity in irradiated metals [32–35,57,58]. It was shown in several fcc and bcc metals that


the true strain hardening behavior is dose-independent[32,35]. The dose-independence of strain hardening wasalso observed in the present study on Mo, as shown inFig. 17, where the true stress-true strain curves for the neu-tron-irradiated Mo are superimposed on the unirradiatedcurve by shifts in the negative or positive direction of thestrain axis depending on the dose and its consequent soft-ening and hardening behavior. It holds true for not onlythe true stress-strain curves tested at different temperaturebut also those tested at varied strain rates. Likewise, theinsensitivity of the PIS to neutron dose was also confirmedin neutron-irradiated Mo. The implication of irradiation-insensitive strain hardening is that one constitutive stress-strain relation may be applied to both unirradiated andirradiated conditions.

Dislocation channels were observed only in the mediumand higher dose specimens (>�0.01 dpa) that exhibitedpremature plastic instability or embrittlement. Dislocationchanneling was not observed in specimens that deformed

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200

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1x10-3 s-1

4

3

2

5


Tru

e st

ress

(M

Pa)

Tru

e st

ress

(M

Pa)

True strain

True strain

1: Unirradiated2: 0.000072 dpa3: 0.00072 dpa4: 0.003 dpa5: 0.0072 dpa

1

0.0 0.1 0.2 0.3 0.40

200

400

600

800

1000Tirr=80oC, Ttest=22oC

1x10-5 s-1

1: Unirradiated2: 0.003 dpa

1 2

Fig. 17. True stress–strain curves for Mo neutron-irradiated at 80 �C andtested at 22 �C (a) strain rate 1 � 10�3 s�1 and (b) strain rate 1 � 10�5 s�1.The curves are superimposed on the unirradiated curve by shifting alongthe strain axis.

uniformly. Dislocation channels were also observed in irra-diated Mo in earlier work [59–61]. However, no definitecorrelation between dislocation channeling and promptplastic instability can be established for irradiated Mo. Itis apparent that the local strain level, test temperatureand the stress state are important factors that affect charac-teristics of dislocation channeling. An undeformed matrixwith a high density of defect clusters and defect-free chan-nels was observed in the 0.0072 dpa/�50 �C specimen at auniform strain of 0.24%, while a highly-deformed matrixwith dislocation networks together with defect-free chan-nels was observed in the 0.072 dpa/22 �C specimen. Ineither case, an increased defect cluster density along thechannel edges was not observed, which rules out the mech-anism of the ‘snow-plough’ removal of defects [62]. Theobservation of channels at �50 �C implies that a diffu-sion-controlled mechanism for defect removal is not impor-tant in irradiated Mo [62]. The reduction in channel widthat a lower deformation temperature suggests that the crossslip of the screw dislocations may be an important factor inthe formation of dislocation channels in Mo [63].

Three fracture modes were observed in neutron-irradi-ated Mo: layered ductile fracture, cleavage fracture andgrain boundary brittle fracture. The fracture mode is moredependent on test temperature and radiation hardeningrather than strain localization. Fig. 18 maps out fracturemodes at a given test temperature and neutron dose in termsof temperature dependence of the PIS, upper yield stress(UPS) (for specimens exhibiting plastic instability at yield)and fracture stress (FS) for Mo. The fracture stress was cal-culated by the final fracture load and the final cross-sectionarea. The fracture mode is most sensitive to test temperature,radiation hardening has secondary importance and strainlocalization has little influence on fracture modes. In fact,premature necking at yield does not seem to affect theamount of necking strain or reduction of area, as shown in

-100 -50 0 50 100 1500

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1400

1600

B

B

B+DD

DB+D

B PIS

Mo irradiated at 80oC

Str

ess

(MP

a)

Temperature (oC)

Unirr 0.000072dpa 0.00072dpa 0.0072dpa 0.072dpa 0.28 dpa

Solid Symbols: PISHalf-open Symbols: YSOpen Symbols: FS

UYS

B

FS

Fig. 18. Plastic instability stress (PIS), upper yield stress (UYS), andfracture stress (FS) as a function of test temperature for unirradiated andirradiated Mo. The fracture modes are marked (D – ductile fracture, B –brittle fracture).

-100 -50 0 50 1000.00

0.05

0.10

0.15

0.20

0.25

0.30Mo neutron-irradiated at ~80oC

Ten

sile

To

ug

hn

ess

(J/m

m3)

Temperature (oC)

0 dpa

0.01 dpa

0.1 dpa

Fig. 19. Tensile toughness vs. test temperature for unirradiated andirradiated Mo. The tensile toughness is defined as the area of the stress–strain curve which corresponds to the dimension of the energy per unitvolume.


Fig. 9(e). The irradiation-induced loss of strain hardenabil-ity may be responsible for the decrease in tensile toughness.As shown in Fig. 19, where the tensile toughness is calculatedby integrating the engineering stress–strain curves, neutronirradiation caused an increase in tensile DBTT and a largereduction in upper shelf tensile toughness that is associatedwith premature plastic instability and ductility loss.

No direct correlation between dislocation channelingand brittle fracture in neutron-irradiated Mo could be con-firmed in these experiments. In fact, the specimen thatexhibited channeling failed in a completely ductile manner,and the specimen that showed extensive uniform deforma-tion showed cleavage fracture. Significant irradiationhardening seems to be a major cause of irradiation embrit-tlement at low temperatures. It has been found that thecleavage fracture stress of a material is insensitive to neu-tron irradiation [64]. The calculated fracture mechanismmap of irradiated Mo has shown that the cleavage fracturefield is enlarged after irradiation due to the insensitivity ofintrinsic cleavage fracture stresses to neutron irradiationand irradiation-induced increase in the yield stress [64].The transition temperature from brittle to ductile fractureshifted from 0.09 Tm for unirradiated Mo to 0.14 Tm forirradiated Mo in the calculated map.

5. Conclusions

The important findings from this study are summarizedas follows:

(1) The major types of irradiation-induced defect struc-ture in Mo were dislocation loops, rafts and cavities.The threshold for visible dislocation loops by TEMwas �0.001 dpa, and the loop density maximumwas at 0.03 dpa. Raft formation became evident at

�0.01 dpa. Cavities were detectable by PAS at alldoses between 7.2 � 10�5 and 0.28 dpa. The meansize of vacancy cavities was about 0.5–0.6 nm. Boththe mean size and size distribution of cavities wererelatively insensitive to neutron dose. The formationof sessile interstitial defect clusters in Mo is suggestedto be dominated by homogeneous nucleation andgrowth. In-cascade vacancy clustering appears to bemore significant in Mo than in bcc Fe.

(2) Neutron irradiation has a pronounced effect on theyield stress of Mo. Irradiation at lower doses reducedthe temperature and strain rate dependence of theyield stress, leading to softening at a lower test tem-perature and at a higher strain rate. The crossoverof softening and hardening is related to the test tem-perature and the strain rate. Further irradiation gaverise to athermal hardening only. An apparent satura-tion or reduced magnitude of irradiation hardeningoccurred at >�0.1 dpa.

(3) The magnitude and the temperature and strain ratedependence of the PIS was apparently not affectedby neutron irradiation. The temperature and strainrate dependence of the yield stress was reduced byirradiation to the same level as of the PIS beforeathermal hardening took place.

(4) Significant irradiation hardening was accompaniedby premature necking at yield and embrittlement.Dislocation channeling was observed in the speci-mens that showed premature necking at yield. How-ever, no direct correlation between dislocationchannels and premature plastic instability can beestablished.

(5) Ductile laminar fracture and transgranular cleavagemixed with intergranular brittle fracture wereobserved in irradiated Mo. Brittle fracture is primar-ily dependent on the test temperature and the strainrate, and the magnitude of irradiation hardening.There is no apparent connection between dislocationchanneling and embrittlement.

Acknowledgements

The research was sponsored by the Office of FusionEnergy Sciences, the US Department of Energy undercontract DE-AC05-00OR22725 with Oak Ridge NationalLaboratory, managed and operated by UT-Battelle, LLC.The authors would like to thank Dr R.E. Stoller for hisvaluable input. We would also like to thank J.L. Bailey,A.M. Williams, L.T. Gibson and P.S. Tedder for their tech-nical support.

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Documents

Low temperature neutron irradiation effects on microstructure and tensile properties of molybdenum