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Temperature dependence of irradiation-induced creep in dilute nanostructured Cu–W alloys

Kaiping Tai a ,?,Robert S.Averback a ,Pascal Bellon a ,Yinon Ashkenazy b ,Brad Stumphy a

a Department of Materials Science and Engineering,University of Illinois at Urbana–Champaign,Urbana,IL 61801,USA b

Racah Institute of Physics,Hebrew University of Jerusalem,Jerusalem 91904,Israel

a r t i c l e i n f o Article history:

Received 10August 2011Accepted 25November 2011

Available online 13December 2011

a b s t r a c t

Irradiation-induced creep (IIC)in dilute Cu–W nanostructured alloy ?lms,300nm thick,was measured as a function of temperature during 1.8-MeV Kr +irradiation using plane-strain bulge testing.The creep rate increased with increasing temperature between 300K and 473K,and then became constant up to 573K.An activation enthalpy of 0.30±0.05eV was obtained for Cu 93.5W 6.5and Cu 99W 1alloys.Thermal creep,in absence of irradiation,became comparable to IIC at 573K.Primary and secondary creep were observed at all temperatures.The steady state creep rate was proportional to the applied stress.Subsequent (scan-ning)transmission electron microscopy analysis revealed a high density of small (2–3nm)W-rich nano-particles with BCC structure after irradiation at all temperatures,but no dislocation loops.The average grain size of the irradiated alloys was stabilized at $30–40nm in both alloys.Correlations between the microstructures and creep behaviors are discussed in terms of grain boundary creep mechanisms.

1.Introduction

Deleterious long-term evolutions of materials under particle irradiation are largely due to the production of sustained net ?uxes of point defects.The origin of these ?uxes is the supersaturation of point defects produced by the irradiation,which is only partially diminished by recombination.Defect ?uxes generally lead to dimensional instabilities in the materials,swelling from the sus-tained growth of vacancy clusters and accelerated creep from the increased rate of climb of dislocations.Defect ?uxes can also result in solute segregation and phase instabilities.In principle,a simple and direct approach to avoiding these microstructural changes is through the addition of a high density of unbiased point defect sinks or traps to the microstructure.While early work explored the use of solute atoms for this purpose [1],more recent ap-proaches have focused on nanostructured materials containing large fractions of incoherent,or semi-coherent,interphase bound-aries [2]or large concentrations of grain boundaries [3–5].An important question regarding these materials,however,concerns their long-term stability under irradiation.Indeed,a vast body of knowledge has been rapidly accumulating about the microstruc-tural stability and strength of nanostructured materials [6–8],but far less is known about their creep stability,particularly under the conditions of high temperatures and intense particle irradia-tion.One reason for this lack of knowledge derives from the dif?culties in performing irradiation-induced creep (IIC)measure-ments,which have mostly been performed on bulk samples in a reactor environment.In past work,we demonstrated that in situ bulge tests on thin ?lms was a convenient means to measure creep rates in nanocrystalline materials using MeV ion irradiation [9].In the present work,we apply this method to measure the tempera-ture dependence of IIC in dilute Cu–W alloys.In addition,we em-ploy (scanning)transmission electron microscopy (TEM)to provide a detailed characterization of the irradiated samples to help eluci-date the underlying creep mechanisms.2.Experimental procedures

The procedures for performing bulge tests on thin ?lm speci-mens in absence of irradiation are described by Vlassak and Nix [10,11].We brie?y summarize these methods,here,emphasizing those parts particular to the current work.A (100)Si wafer was ?rst coated on both sides with a Si 3N 4?lm,$300nm thick,using plasma-enhanced chemical vapor deposited (PECVD).The Cu–W ?lms were subsequently deposited on the wafer using DC magne-tron sputtering with individual Cu and W targets (99.999%).The base pressure in the growth system was $2.67?10à6Pa,and the operating Ar pressure during growth was %0.267Pa.Rutherford backscattering was employed to measure the thick-nesses of the ?lms,$300nm,and the composition of Cu–W alloys,Cu 93.5W 6.5and Cu 99W 1.A layer of polymer bond was coated onto the Cu 93.5W 6.5/Cu 99W 1?lms to protect them during subsequent processing.Photolithography was next used to de?ne long rectan-gular windows in the PECVD Si 3N 4coating on the backside of the substrate.The Si substrate was then etched anisotropically using

Corresponding author.

E-mail address:kptai@https://www.wendangku.net/doc/713019440.html, (K.Tai).

a hot potassium hydroxide (KOH)solution to membranes consisting of the Si 3N 4coating ?lms.The size of the window membranes is ly,the freestanding Cu 93.5W 6.5/Cu 99W 1removing the Si 3N 4using reactive ion ing the protective polymer layer in an mens were then sealed to the sample vacuum (and irradiation)on one side of the and gas pressure on the other.

The sample was heated using tungsten symmetrically around a ceramic pillar on the ple holder.The temperature was measured couple,located very close to the sample.the sample holder ensured good temperature sample.The measured temperature stability The samples were irradiated with 1.8MeV Kr tween $5?1019to $1.5?1020Kr +/m 2(%rates were %4?10à3dpa s à1.Pressure was 7%H 2mixture of gas,the H 2preventing at elevated temperature.The de?ection or of the membrane was measured using a a displacement resolution of 316.5nm,i.e.,the He–Ne laser.It should be noted that the ?lm specimens,as will be shown below,is %less than the ?lm thickness.Surfaces effects be signi?cant in these experiments.

Analysis of the bulge testing was taken rectangular freestanding membrane with the transverse component of the biaxial of gas pressure and membrane de?ection is r x

p ea 2th 2

T2ht

where p is the applied gas pressure,t is the the bulge height in the middle of the sample,of the membrane.The corresponding strain dent of whether the ?lm is deformed be calculated as [11]:

e x

?e 0t2h 2

3a 2arcsin 2ah

a 2th

2 à1where e 0is the residual plane strain in the a scanning laser beam method,which is more curvature than is laser interferometry,was the initial strain in the sample at the different The Cu–W samples were characterized testing using transmission electron ning transmission electron microscopy ?rst milled to less than %30nm in thickness W has a large atomic number relative to Cu,high angle annular dark ?eld (HAADF)imaging (Z -contrast)could be used to analyze the nano-sized precipitates.3.Experimental results 3.1.Irradiation creep

The irradiation creep curves of Cu 93.5W 6.5and Cu 99W 1alloy ?lms are plotted in Figs.1and 2as plane strain versus ?uence (in dpa)for various temperatures.For some of the creep tests,the applied pressure was increased during the irradiation in dis-crete increments as marked in the two ?gures.All of the specimens are observed to undergo primary creep for a short time (dose)before reaching a near steady state regime.Primary creep is also observed after each change in the applied stress during the

irradiation.The duration of the primary creep is %5dpa,although this dose appears to decrease somewhat with additional stress increments later in the irradiation.Creep measurements were con-tinued after the ion beam was switched off to evaluate the contri-bution of thermal creep.The microstructures of the samples at this time are presumably similar to those at the ?nal stages of the IIC measurements.The thermal creep rate after IIC at 473K and an ap-plied pressure of 1KPa was $6.6?10à8s à1,which represents a factor of $1/6th the total measured IIC.After the IIC measurements at 573K,the thermal creep increased by a factor of $4over that at 473K,and now accounting for $1/2of the measured IIC.Thermal creep measurements were also performed at 473K and 573K on samples without prior irradiation;these data are plotted in Fig.3.The steady state creep rates at 473K were 1.06?10à7s à1and 1.59?10à7s à1for applied pressures of 1KPa and 1.5KPa

Fig.1.Creep plane strain-dpa curve for Cu 93.5W 6.5freestanding ?lms irradiated by 1.8MeV Kr +at 300–573K.For clarity,each curve is successively offset 0.001for the data from ()to (h ).Notes:(h )IIC at 573K &1.5KPa,()thermal creep at 573K 1.5KPa,()IIC at 573K &1.4KPa,()thermal creep at 573K &1.4KPa,()IIC 573K &1KPa,()IIC at 473K &1,2,3KPa,()IIC at 373K &1KPa,()IIC 348K &1KPa,()IIC at 300K &1KPa.Black arrows indicate the stages of applied gas pressure at 473K (A:1KPa,B:2KPa and C:3KPa).Irradiation rate:$0.004dpa/For thermal creep time is converted to dpa using the displacement rate.

Fig.2.Creep plane strain-dpa curve for Cu 99W 1freestanding ?lms irradiated by 1.8MeV Kr +at 300–473K.Notes :()IIC at 473K &1,2,3kPa,()thermal creep 473K &3KPa,()IIC at 373K &1KPa,(j )IIC at 328K &1KPa,()IIC at 300K KPa.Black arrows indicate the stages of applied gas pressure at 473K (A:1KPa,KPa and C:3KPa).Irradiation rate:$0.004dpa/s.For thermal creep time converted to dpa using the displacement rate.

and thus nearly linear in stress.At 573K the creep rate increased slightly to 1.28?10à7s à1for 1KPa pressure.This small change in creep rate with increasing temperature appears to be due to a dif-ference in the microstructures of these unirradiated samples,but this will be discussed later.

The effect of applied pressure on irradiation-induced creep rate was obtained from the creep data shown in Figs.1and 2.These re-sults are reported in Fig.4,where the steady state creep rates are plotted as a function of von Mises equivalent stress for both the Cu 93.5W 6.5and Cu 99W 1specimens.The data are well ?t by a linear dependence on stress,with the strain rate extrapolating to zero at

zero stress.The creep compliance,B ?_e

=r e ,can thus be obtained from slopes of these lines.

The temperature dependencies of IIC for the two Cu–W alloys

being a factor of %20times larger than that at room temperature.At 573K,however,the normalized creep rates do not change from those at 473K,suggesting the IIC becomes constant above %473K in these two alloys.A best ?t to the data between room temperature and 473K yields an apparent activation energy of %0.30±0.05eV.

3.2.Microstructural analysis

The microstructures of the Cu 93.5W 6.5and Cu 99W 1?lms were analyzed using TEM and STEM.A TEM micrograph of Cu 93.5W 6.5?lm in the as-deposited state,before IIC deformation,is shown in Fig.6.The average grain size of these samples,which have moder-ate densities of twins,was $30–40nm.High angle annular dark Fig.3.Creep plane strain–time curve in Cu 93.5W 6.5freestanding ?lms.Notes :(j thermal creep at 473K &1KPa,()thermal creep at 473K &1.5KPa,()thermal creep at 573K &1KPa.

Fig.4.Dependence of strain rate/dose rate on the (von Mises)equivalent uniaxial stress in Cu 93.5W 6.5and Cu 99W 1.Creep compliance,B ,is provided.

Fig.5.Normalized creep rate,kT e RED =r e ,as a function of 1/T for Cu 93.5W 6.5and

Cu 99W 1.

Fig.6.TEM bright ?eld image of Cu 93.5W 6.5as-deposited state.

the Z-contrast STEM images)are uniformly distributed in the grain interiors and grain boundaries.Precipitates in the grain inte-riors are found by diffraction to be(semi-)coherent with the matrix.An example is shown in the high resolution TEM Fig.8b.The geometry of the W precipitates at grain boundaries oval,and within the grain interior it is spherical.Energy dispersive spectrometry analysis con?rmed that the small($3nm)white spots in these?gures are indeed W-rich nanoparticles.Similar pre-cipitate formation during IIC was observed at the other irradiation temperatures as well.The density of W precipitates in Cu99W1 loy was signi?cantly lower than in the Cu93.5W6.5,by a factor four,but the size of the particles was about the same,smaller than nm(see Fig.9).These general observations agree well with pre-vious?nding using X-ray diffraction that W precipitates out solution in Cu90W10under ion irradiation,even at room tempera-ture,but not during thermal annealing at temperatures up to

800K[14].MD simulations in Ref.[14]suggest that the W precip-itates form within the thermal spike phase of the cascade evolution.

4.Discussion

The principal?ndings in this work can be summarized as follows:

IIC rates in nanocrystalline Cu93.5W6.5and Cu99W1are the same, within experiments uncertainties,at all temperatures examined.

The grain sizes of these two alloys are also nearly the same.

W precipitates are formed during irradiation in both alloys,but not in the unirradiated alloys.The size of the precipitates,which are observed both in grain interiors and grain boundaries,are

smaller than$3nm,with the densities(in atom fraction)of %2.4?10à5and%5.9?10à6in the Cu93.5W6.5and Cu99W1 alloys,respectively.

No evidence for dislocations in the grain interiors is observed in the irradiated samples.Dislocations were found,however,after irradiation of undoped Cu samples with larger grain sizes, %100nm[9].

IIC depends on temperature over the range300–473K,and then becomes constant up to573K.At this temperature,thermal creep becomes signi?cant.The apparent activation of creep in the temperature dependent regime is%0.3eV.

The key observation in this study is that the temperature dependence of IIC follows closely that found for radiation en-hanced diffusion(RED)in Cu[15].At low temperatures,RED is in the so-called recombination regime.The diffusion coef?cient is

Fig.7.TEM bright?eld image(a)and HAADF STEM image(b)of Cu93.5W6.5after

thermal creep deformation at573K.

Fig.8.TEM bright?eld image(a),HRTEM image(b)and HAADF STEM image(c)

93.5

W6.5after IIC deformation at573K.

temperature dependent with an apparent activation energy of one-half the migration energy of vacancies,D H m m =2.For Cu,D H m

m ?0:71eV [16],which is approximately twice the value observed here for IIC.At higher

temperatures RED enters the sink limited regime,where RED becomes independent of temperature.At still higher temperatures,thermal diffusion becomes dominant.These general characteristics of RED are all found in the IIC data well.

Since the TEM observations fail to show any evidence for dislo-cations forming during irradiation of these alloys,we attribute the creep behavior to point defects ?owing to grain boundaries.We defer a detailed discussion of how defect annihilation at GB’s gives rise to creep to a later publication on modeling of the creep process [17],but we mention here that (i)the creep rate in this model kinetically limited by the arrival of point defects at the GB’s and (ii)it varies as 1/d rather than 1/d 3characteristic of Coble creep.The IIC values reported here,therefore,should not be overly sensi-tive to our measurement of the grain size,which have an uncer-tainty of %20%.The model further suggests that the temperature dependence of IIC should be given by the temperature dependence of the number of point defects annihilating at grain boundaries.We next estimate this temperature dependence alloys.

From chemical rate theory we write [18]@c i em T@t ?n FM K 0à

4p r i m X eD i tD m Tc i em Tc m ei Tà4p r p

X à3p 2

n gb D i ;m c i em T

where c i (v )is the concentration of fraction,X is the atomic volume,r iv and r p tial-vacancy recombination and grain size,n p (gb )is the sink ef?ciency of precipitates (gb’s),n FM is the fraction of freely migrating point defects,i.e.,point defects that escape the cascade [19,20],and D i (v )is the diffusivity of the intersti-tial (vacancy)https://www.wendangku.net/doc/713019440.html,parison of the second term with the two sink terms determines whether the system is in the recombination or sink dominated regime,while comparison of the third and fourth terms yields the fraction of defects annihilating at gb’s compared to the fraction annihilating at precipitates.

We ?rst calculate the ratio R of defects annihilating at gb’s ver-sus https://www.wendangku.net/doc/713019440.html,ing the values,c p =2.4?10à5,r p =1.5nm,L =30nm for the Cu 93.5W 6.5alloy,and c p =5.9?10à6,r p =1.5nm,L =30nm for the Cu 99W 1alloy,we obtain R %3f gb =f p and R %0:9f gb =f p for the two alloys,respectively.Thus,if the sink ef?-ciencies of precipitates and GB’s are similar,the sink strengths are similar.We therefore neglect the precipitate sinks in what follows since a factor of two in the analysis has little bearing on the conclu-sions.We point out,however,that the W precipitates are semi-coherent in the Cu matrix,and hence their sink ef?ciency may be smaller than those of high energy GB’s [7].The result that the two Cu–W alloys have very similar values of IIC perhaps supports this hypothesis,since their precipitate densities differ by a factor of %4.The uncertainties in the grain size,grain shape and precipitate density in the two alloys,however,preclude a de?nite conclusion at this time.

The ?ux of defects into the GB’s can be obtained from Eq.(3)using D m ?exp eà0:71=kT T,r iv =a 0,and K 0=4?10à3dpa s à1,and L =30nm.We assume n FM is between 0.02and 0.10[19,20].The temperature dependence of the ?ux of defects into the GB’s is shown in Fig.10in an Arrhenius plot.The data points in this plot show the actual creep data.The agreement between the tempera-ture dependencies of the calculated curve and the experimental data is remarkably good,illustrating that between 300K and 473K,creep is occurring in the recombination regime,and above this temperature,the system enters the sink-limited regime.This analysis,coupled with the failure to observe dislocations with the grain interiors,thus provides the ?rst strong evidence for a GB mechanism of IIC.

Fig.9.TEM bright ?eld image (a)and HAADF STEM image (b)of Cu 99W 1after irradiation creep deformation at 473K.

Fig.10.Calculated GB sink effects and experimentally measured creep rate for Cu 93.5W 6.5and Cu 99W 1.Notes :The values are normalized at 473K.

Materials 422(2012)8–13

where d D GB is the product of the GB width and GB diffusion coef?https://www.wendangku.net/doc/713019440.html,ing diffusion data from Ref.[22],we obtain_e%1?10à3 at500K,or about4orders of magnitude larger than the measured value.While we can only speculate about this difference here,we note that solute segregation to grain boundaries can strongly affect GB diffusion.MD simulations at much higher temperatures, %1050K,indeed show oversized solutes in Cu(Nb and Zr)strongly suppress thermal creep[23],and W is oversized in Cu as well.The weak temperature dependence in the creep rate between473K and 573K is also unusual.In this case,however,the microstructure is changing with temperature.For example,the grain size increased from$30nm to$50nm between473K and573K,which accord-ing to Eq.(4)would decrease the creep rate by a factor of%5.

5.Conclusions

We have shown that creep deformation in nanocrystalline Cu–W alloys is greatly enhanced by ion irradiation in the temper-ature range300K to573K.Analysis of the temperature dependence of the creep rate suggests that IIC in these ultra?ne grain alloys is controlled by point defect?uxes to the grain bound-aries.Below473K,the alloy is in the so-called recombination regime,where point defects predominantly annihilate by recombi-nation reactions.By573K,most of the defects annihilate at either grain boundaries or at precipitates,although,comparison between the Cu93.5W6.5and Cu99W1alloys suggests that the(semi-)coherent W precipitates are not effective sinks for point defects.These measurements thus reveal,for the?rst time,a GB mechanism of irradiation creep.The thermal creep rate in these alloys,in absence of irradiation,is surprisingly small,being four orders of magnitude smaller than that predicted by Coble creep.We tentatively attribute this small creep rate to lowering of the grain boundary energies by the presence of W in the boundaries,although additional work will be necessary to understand this behavior.

This research was supported by the US DOE-BES under grant DEFG02-05ER46217.It was carried out,in part,in the Frederick Seitz Materials Research Laboratory Central Facilities,University of Illinois,which are partially supported by the US Department of Energy under grants DE-FG02-07ER46453and DE-FG02-07ER 46471.

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