effect of ta solute concentration on the microstructural evolution...

Effect of Ta Solute Concentration on the MicrostructuralEvolution in Immiscible Cu-Ta Alloys

B.C. HORNBUCKLE,1,5 T. ROJHIRUNSAKOOL,2 M. RAJAGOPALAN,3

T. ALAM,2 G.P. PURJA PUN,4 R. BANERJEE,2 K.N. SOLANKI,3

Y. MISHIN,4 L.J. KECSKES,1 and K.A. DARLING1

1.—Weapons and Materials Research Directorate, Aberdeen Proving Ground, Maryland 21005,USA. 2.—Center for Advanced Research and Technology, Department of Materials Science andEngineering, University of North Texas, Denton, TX 76203, USA. 3.—School of Engineering ofMatter, Transport, and Energy, Arizona State University, Tempe, AZ 85281, USA. 4.—Depart-ment of Physics and Astronomy, George Mason University, Fairfax, VA 22030, USA. 5.—e-mail:[email protected]

The immiscible Cu-Ta system has garnered recent interest due to observationsof high strength and thermal stability attributed to the formation of Ta-en-riched particles. This work investigated a metastable Cu-1 at.% Ta solidsolution produced via mechanical alloying followed by subsequent consolida-tion into a bulk specimen using equal channel angular extrusion at 973 K(700!C). Microstructural characterization revealed a decreased number den-sity of Ta clusters, but with an equivalent particle size compared to a previ-ously studied Cu-10 at.% Ta alloy. Molecular dynamic stimulations wereperformed to understand the thermal evolution of the Ta clusters. The clustersize distributions generated from the simulations were in good agreementwith the experimental microstructure.

INTRODUCTION

Nanocrystalline materials have, and continue toreceive, significant interest for their ability toenhance mechanical properties, specifically,strengthening via the Hall–Petch mechanism,compared to materials with a coarse-grainedmicrostructure. While the nanometer-sized grainsare clearly advantageous in imparting greaterstrength to the material, they do not contribute to,or frequently are detrimental to, the plasticity of thematerial. Furthermore, these nanometer-sizedgrains are highly susceptible to grain growth due totheir thermal instability, which limits their use inpractical applications. To address this limitation,considerable research has focused on enhancing thethermal stability of such systems to prevent/mini-mize the onset of grain growth. As a result, twomethods have come to the forefront: a kineticapproach and a thermodynamic approach. The for-mer utilizes several mechanisms to apply a pinningforce on the grain boundaries (GBs) to prevent graingrowth, while the latter suppresses grain growth byreducing the GB free energy through solute atoms

partitioning to the GBs.1–7 The mechanisms typi-cally employed to pin the GBs are secondary pha-ses8–11 (Zener pinning) and/or solute drageffects.12,13

In light of these two competing mechanisms,there has been much discussion on which methodmay provide a more successful path in bringingabout the realization of commercially available bulknanocrystalline metals. In many cases, such debatehas been fostered by the fact that it is not alwayspossible to fully separate or delineate the contribu-tions of these two competing stabilization mecha-nisms in preventing grain growth in nanocrystallinemetals. For instance, thermodynamic stabilizationof nanocrystalline grain size (GS) involves examin-ing the energetic penalty associated with the highvolume fraction of the GBs, and the possibility ofsolute segregation driving this associated excessfree energy to zero.14–17 However, intertwined inthis scenario are the kinetic aspects of solute drag,and its role in reducing grain growth in this ther-modynamic stabilization construct, which has beenan area of active research.18,19 Additionally, theprecipitation of secondary, solute-rich phases has

JOM, Vol. 67, No. 12, 2015

DOI: 10.1007/s11837-015-1643-x" 2015 The Minerals, Metals & Materials Society (outside the U.S.)

2802 (Published online September 28, 2015)

http://crossmark.crossref.org/dialog/?doi=10.1007/s11837-015-1643-x&domain=pdf

http://crossmark.crossref.org/dialog/?doi=10.1007/s11837-015-1643-x&domain=pdf

been experimentally observed to disrupt the stabi-lization set in place by the thermodynamic mecha-nism. Nevertheless, recent research has shown thatjust because phase separation or precipitationoccurs does not necessarily mean that a stabilizednanocrystalline system does not exist. That is,recent theoretical work by Murdoch and Schuh haspredicted the existence of stable duplex systems,wherein both GB segregation and phase separationoccurs, resulting in both a stable nanocrystalline GS(i.e., GB energy of zero) and the precipitate struc-ture coexisting with one another.20 These types ofmicrostructures are currently under investigation.Still further, significant stabilization by the Zenerpinning of GBs has been experimentallyreported.14,21

The convolution and competition of these kineticand thermodynamic effects, and the complex rela-tionships between their alternative stabilizingmechanisms in nanocrystalline alloys, continues toremain an important aspect of research in the nearand distant future. One particular class of systems,immiscible Cu-based alloys, has had the attention ofthe research community due to their outstandingmechanical properties and resistance tomicrostructural coarsening at elevated tempera-tures.22 In general, these systems are intriguing, asthey posses attributes which suggest that bothmodes of stabilization, i.e., kinetic and thermody-namic, may play a role in stabilizing theirmicrostructures, with Cu-Ta being one of thesesystems.22–28

Recently, it has been shown that nanocrystallineCu-10 at.% Ta can retain a mean GS of 167 nm evenafter annealing at 97% of its melting point (TM).

25

This extreme increase in microstructural stabilitywas suggested to be a combination of thermody-namic and kinetic stabilization effects which, inturn, appeared to be controlled by segregation anddiffusion of Ta solute atoms along GBs. The forma-tion and precipitation of secondary phases compli-cated the ability to decipher the underlyingcontrolling stabilization mechanism, but it wassuggested that, at moderately high temperatures,the microstructure is stabilized primarily by kineticpinning effects and only slightly, if at all, by ther-modynamic factors. To help provide greater clarityto the situation, computational work was performedto understand the behavior of Ta when distributedin Cu along GBs or dispersed as a random solidsolution. The work indicated that a GB-segregatedsolute was significantly more effective at preventinggrain growth; however, at higher temperatures, itwas prone to short-circuit diffusion, leading to theformation of Ta clusters along the GBs.23 Despitethe higher temperatures, these clusters effectivelyprevented further grain growth, and pointed to thesignificance of Zener pinning and the importance ofTa clusters in preventing grain growth in thissystem.

In the present work, new atomic simulations,which specifically investigate the sole effect of Taclusters, preventing grain growth in nanocrystallineCu,29 are contextualized in relation to new atomprobe tomography (APT) data. These computationalsimulation results are compared and found to be inagreement with the experimental results onnanocrystalline Cu with 1 at.% and 10 at.% Taprocessed at 973 K (700!C). The role and evolutionof composition, number density, and cluster size arefound to be interrelated. These aspects are thendiscussed in relation to grain growth and mechani-cal properties. These new findings point to thedominant stabilizing mechanism being kineticrather than thermodynamic in nature. and offer amore refined methodology for developing stable bulknanocrystalline alloys, which retain their extraor-dinary properties post-processing and duringapplication.

Experimental Procedures

The ability to produce this immiscible systemwith 5 at.% or greater of Ta has been limited to afew select processes: ion beam mixing, physicalvapor deposition, or ball milling. High-energy ballmilling has been utilized specifically for the twocompositions, Cu-1 at.% Ta and Cu-10 at.% Ta,produced in this study. For each of these two com-positions, appropriate amounts of Cu and Ta pow-ders (!325 mesh, 99.9% purity) with total weight of5 g were loaded into hardened steel vials along withthe milling media (440C stainless steel balls) with aball-to-powder ratio of 10-to-1 by weight. Theweighing out of the powder and milling media plusthe sealing of the vials were performed in a glovebox under an Argon atmosphere (O and H2O are<1 ppm). The milling process was performed atcryogenic temperature (verified to be 77 K("!196!C)) in a SPEX 8000 M shaker mill. Toensure that the vial reached the desired tempera-ture, it was allowed to equilibrate for at least 1200 s(20 min) prior to milling. Both powders were milledfor a duration of 28.8 ks (8 h). Once the millingprocess was completed, the vials were placed backinside the glove box before opening, and the powderremained there until further processing. This cryo-genic milling process was repeated until 100 g ofeach composition was produced. The resultingpowder was an unagglomerated mass with a par-ticulate size ranging from 20 to 100 lm.

For consolidation of the powders, equal channelangular extrusion (ECAE) was selected. For theECAE consolidation process, the powders wereplaced into nickel cans and sealed within the glovebox. The nickel cans’ dimensions were 90 mm inlength with a cross-section of 25.4 mm 9 25.4 mmwith a "10-mm-diameter hole (with a length of"50 mm) bored along the long axis of the cans. Onceloaded with the proper powder amount, the cans

Effect of Ta Solute Concentration on the Microstructural Evolution in Immiscible Cu-Ta Alloys 2803

were heated to 973 K (700!C) for 2400 s (40 min)under an Ar atmosphere to ensure that the powdermass was thermally equilibrated. To minimize heatloss between the die and can, the die assembly wasalso preheated but to a temperature of 573 K(300!C). Once equilibrated, the cans were quicklydropped into the ECAE tooling, where they wereextruded at a rate of 25.5 mm s!1. To prevent/min-imize texture effects, the cans were extruded fourconsecutive times following route Bc.

30–34 With theECAE tooling having a 90! channel angle, the fourconsecutive extrusions resulted in a total strain of"450%. After extrusion, the powder consolidate wasremoved from the can using a wire electric dis-charge machining. No porosity or particle bound-aries from the initial powders were found whenobserving the consolidated powder using a scanningelectron microscope.

Microstructural analysis was performed using aJEOL ARM 200F transmission electron microscope(TEM) operated at 200 kV. TEM specimens wereprepared from the 1 at.% and 10 at.% Ta powdersECAE processed at 973 K (700!C) both in a con-ventional manner and using the lift-out method.The conventional specimens were made by punch-ing out 3-mm-diameter disks from slices of theconsolidated powder with a Gatan 659 disk punch.The disks were first thinned to a thickness of"50 lm. The specimen was then dimple-grinded toleave a final dimple thickness of £5 lm. Finally, thespecimen was Ar ion-milled until electron trans-parent in the dimple. The lifted-out specimens weremade using a FEI Nova Nano Lab 600 focused ionbeam (FIB) unit.

The atom probe tomography (APT) analysis wasperformed using a CAMECA LEAP 3000XTM HRinstrument. The specimens were lifted-out to gen-erate 8 tips and then annular-milled to a final tipdiameter of 50–70 nm using the FEI Nova200TM

dual-beam focused ion beam system. The specimenwere held at a base temperature of 45 K (!228!C),while the experiment was conducted in the pulsed-laser evaporation mode, with a laser pulse energy of0.3 nJ, and target evaporation rate of 0.5%. TheAPT data was reconstructed and quantitative eval-uation completed using the CAMECA IntegratedVisualization and Analysis Software (IVAS) 3.6.8TM

software. Additionally, at least three individualsample reconstructions were analyzed for eachcondition to ensure microstructural homogeneity.

RESULTS

TEM of Cu-1Ta and Cu-10Ta

Figure 1 is a collection of TEM bright field andscanning transmission electron microscopy high-angle annular dark field (STEM-HAADF) imagesfrom both Cu-1Ta and Cu-10Ta ECAE processed at973 K (700!C). The bright field images (Fig. 1a andc) confirm that the 10% Ta alloy has a smaller GScompared to the 1% Ta alloy, 70 nm and 168 nm,

respectively.24 Both STEM-HAADF images (Fig. 1band d) show the presence of Ta particles throughoutthe respective microstructures. It is apparent thatthere is no set particle size, but rather a range ofparticle sizes within these alloys. While both alloyscontain these Ta-rich particles, the 10% Ta alloyclearly has a higher number of smaller particles(diameter<10 nm), as well as much larger particles(diameter "50 nm to 100 nm) present as comparedto Cu-1Ta. These smaller particles have been sug-gested to be the primary reason for the yieldstrength of this alloy being "1100 MPa in quasi-static compression, which is much greater than thatof pure Cu with the same GS.24 It is also a factor of2–3 times greater than that provided by Hall–Petchbased strengthening alone. It can be further shownthat same level of strength attained in pure Cu, fora GS of 5 nm, can be attained in Cu-Ta alloys withmuch larger GSs ("70 nm to 200 nm). These resultssignify a paradigm shift in alloy development. Theyindicate that it is possible to engineer alloys withlarger grains, i.e., better ductility/plasticity, whilemaintaining strength levels equivalent to or greaterthan that provided by reducing GS to the limits ofnanocrystallinity.

Morphology and Number Density of Cu-1Tafrom APT

From the TEM micrographs, it is visually obviousthat the Cu-10Ta alloy has a higher number of Ta-rich particles present than the Cu-1Ta alloy, butmore quantitative data are required for both thecomposition and number of Ta-rich particles presentwithin these alloys before a full understanding isgained. To achieve this objective, APT was per-formed to observe any compositional differences andmeasure the quantitative number density for theparticles/clusters of these alloys.

Figure 2a provides a three-dimensional (3D)reconstruction of the ECAE-processed Cu-1Tasample found to have a measured bulk compositioncontaining 0.89 at.% of Ta, which is only slightlybelow the global concentration of the alloy. Thereconstruction shows random distribution of largeTa particles (shown in solid black) within the Cumatrix (shown in orange). These large Ta particleswere delineated by using a 5 at.% Ta isoconcentra-tion surface (isosurface). These spheroidal Ta par-ticles have a size ranging from 5 to 30 nm. Apartfrom the large Ta particles, Ta solute atoms (shownin black) in solid solution are also observed withinthe Cu matrix atoms (shown in orange), as shown inFig. 2b.

After removing the large Ta particles from thereconstruction, it was found that 0.54 at.% of Taremains in solution within the Cu matrix. To verifythe existence of very fine Ta-enriched particleswithin the Cu grains, cluster analyses of APTdatasets were employed. The Cu matrix was ana-lyzed using a maximum separation algorithm for

Hornbuckle, Rojhirunsakool, Rajagopalan, Alam, Purja Pun, Banerjee,Solanki, Mishin, Kecskes, and Darling

2804

cluster identification,35–37 within the framework ofthe IVAS 3.6.8 software. This search was first con-ducted on particles enriched in Ta. Ta ions weregrouped together within a maximum (first nearestneighbor) separation distance of 1.15 nm (dmax); agroup of Ta ions with more than 30 ions (Nmin) wasidentified as a cluster. All Cu ions within an envel-ope distance (L) of 0.55 nm of these Ta ions wereincluded as part of that cluster. However, any ionlocated at a distance greater than 0.55 nm (derosion)of any other ion along the perimeter of this desig-nated cluster was excluded from the cluster.

Figure 2c presents Ta-enriched clusters insidethe Cu grains (atoms are color-coded as previouslyshown). Local Ta clusters were found to account fora portion of Ta solid solution ("0.12 at.% Ta) withthe remaining amount of the solid solution("0.42 at.% Ta) composed of randomly distributedTa atoms. The Ta cluster size distributions in theCu matrix for the APT experiments are presented inFig. 2d. Cluster size was adjusted to reflect the trueparticle size due to local magnification effects byusing density corrected composition profiles. The

details of the corrections will be discussed in a latersection. The average Ta cluster size is 3.4 nm (s-tandard deviation = 1.3) prior to the density cor-rection procedure and the number density of Taclusters is 4.1 9 1023 m!3. This number density ofTa clusters is "37% less than the value reported,6.5 9 1023 m!3, for the Cu-10Ta alloy.26

Due to a strong difference in atomic densitybetween Cu-rich and Ta-rich regions in the APTreconstructions, artifacts such as aberrations of iontrajectories and local magnification are likely tooccur.38–41 These artifacts influence true particlesize and the composition in their vicinities.38,39 Forexample, in the Cu-Ta system, if a protrusion at thetip surface develops, a higher magnification andlower atomic density within the high evaporationfield can result, especially for Ta particles, eB = 1.47(44/30).28 Therefore, the apparent sizes of large andsmall Ta particles (20 nm and 4 nm, respectively),were overestimated. Gault et al.40 proposed that thedistance units, in composition profiles containing nbins, need to be adjusted proportionally to thesquare root of the atomic density. Thus, the

Fig. 1. (a) TEM bright field and (b) STEM-HAADF image showing the microstructure of Cu-1Ta ECAE processed at 973 K (700!C), (c) TEMbright field and (d) STEM-HAADF image showing the microstructure of Cu-10Ta ECAE processed at 973 K (700!C).


corrected distance, xi can be adjusted as,

xi ¼ xmax $Pi

j¼1

ffiffiffiffidj

pPn

j¼1

ffiffiffiffidj

p where xmax is the maximum

distance in the composition profile and dj is theatomic density at the jth bin.

Note that the precipitate size obtained from den-sity corrected composition profile (not shown here)is an estimate and can be used for qualitative com-parison. With the Gault’s corrections, the averagesize of the Ta particles decreased from 3.4 nm to1.3 nm. When applying the same correction to theCu-10Ta alloy, the sizes of the Ta particles were alsofound to decrease from 3.4 nm to 1.5 nm.26 Thus, wearrive at the remarkable result that the cluster sizeis nearly identical even though the bulk Ta contentis an order of magnitude different between the twoalloys. In other words, the cluster size is virtuallyindependent of the bulk Ta content. The increasedbulk Ta content results in a higher number densityof particles/clusters rather than yielding enlargedparticle/cluster sizes.

Molecular Dynamics Simulations of Pure Cuand Cu-1Ta

Molecular dynamics (MD) simulations of pure Cuand Cu-1Ta were conducted to better understandhow increased temperature affects Ta and what rolethese factors play in controlling grain growth. Thesimulation of pure Cu clearly displayed significantgrain growth between the non-annealed andannealed states at the temperature of 1200 K (927!C)(see Fig. 3a and b). When comparing the simulationsof pure Cu and the Cu-1Ta alloy under the sameannealing conditions (Fig. 3c and d), clustering of Tawas observed in the alloys already in the pre-an-nealed state. The annealing yielded coarsening ofthese clusters, especially along the GBs. The clustersize distribution plot for the Cu-1Ta alloy (Fig. 4)showed that the most probable cluster size onlysomewhat coarsens from 0.7 nm to 0.8 nm. The veryexistence of this peak and its shift suggest that thesenanoclusters constitute a well-defined and repro-ducible microstructural feature of the alloy.

Fig. 2. Cu-1Ta ECAE processed at 973 K (700!C): (a) APT reconstruction showing large Ta particles, delineated using a 5 at.% Ta isocon-centration surface, in Cu matrix, (b) solid solution of Ta within Cu matrix, (c) APT reconstruction showing Ta-enriched cluster present in the Cumatrix, and (d) cluster size distribution.


2806

As such, it is hypothesized that the solid solutionobtained by ball milling is thermodynamically unsta-ble and is likely to decompose by spinodal decompo-sition resulting in the formation of coherent Taclusters inside the grains. The growth of such intra-grain clusters continues until the increasingly largerelastic strain energy overtakes the chemical drivingforce and frustrates continued growth. Conversely, ifsuch clusters nucleate at GBs, dislocations, and other

defects, it is possible that the growth can continue to amuch larger extent via Ostwald ripening kineticallycontrolled by short-circuit at GBs or pipe diffusionalong dislocations. By a large extent, these observa-tions have already been reported in experiments thatshow preferential coarsening of precipitates at GBs,relative to those in theCumatrix. Note, however, thatthe rate and overall extent of coarsening of theseparticles, as seen experimentally in Fig. 1, is depen-dent on the bulk Ta content. Due to limitations in theavailable time scale, the kinetics and limits of suchcoarsening could not be fully captured by the simula-tions. Nevertheless, the simulations do show correctlythat the vast majority of the Ta clusters form at GBsand triple junctions, and that these clusters play asignificant role in preventing grain growth throughZener pinning.

Still further agreement between simulations andthe experimental APT results can be gained byreferring back to the cluster size distribution plot(Fig. 4). The presence of clusters<0.5 nm indiameterindicates that, besides these clusters present in themicrostructure, there are a small amount of single Taatomsand small groups of 2–3 atomclusters scatteredthroughout the microstructure. The coexistence ofthese twomicrostructural features (clusters and solidsolution) is consistent with current APT results forboth 1 at.% and 10 at.% Ta.26 Size distribution plotsfrom atomic simulations29 (not shown here) also

Fig. 3. Nanocrystalline Cu and Cu-1at% Ta alloy (a, c) before and (b, d) after a 10 ns isothermal MD simulation at 1200 K (927!C). The GBs arerevealed by bond angle analysis using OVITO.42 The Ta atoms are shown in green and are larger than the Cu atoms.

Fig. 4. The cluster size distribution in Cu-1Ta alloy before and after a10 ns anneal at 1200 K (927!C).


suggest that the cluster size remains nearly the samewith increases in composition, and that, with higherTa concentrations, a higher number density of clus-ters is observed. This is also consistent with experi-mental results reported here.

Compositional Examination of Ta Clusters inCu-1Ta and Cu-10Ta Alloys

During high energy mechanical alloying, it isimportant to realize that some level of oxygen can beincorporated into the nanocrystalline lattice, due tothe presence of the native oxide in the startingpowders used. Due to the APT’s ability to detecteven trace elements, it was utilized to observe thebehavior of oxygen in regard to the overall bulkcomposition of the alloy and the clusters them-selves. Low levels (0.05–1 at.%) of Fe contaminationwere already previously reported in similar sam-ples.26 In the same manner as reported for the Cu-10Ta alloy,26 the bulk composition of the Cu-1Taalloy contained 1.21 ± 0.03 at.% oxygen, ("0.25wt.%). Both these two alloys were ECAE processedat 973 K (700!C). In preparation for the ECAEprocessing, both powder compositions were sealedin a Ni can inside a glove box with an Argon atmo-sphere (O and H2O are <1 ppm) and, therefore,exposure to oxygen was minimal during handlingand extrusion. As a precautionary measure, thebulk oxygen content was determined for the Cu-10Ta samples in the as-milled and annealed informing gas at 723 K (450!C) conditions. Thesesamples contained 1.18 ± 0.08 at.% and 1.01 ±0.04 at.% oxygen, respectively. Since the bulk oxy-gen content of all four samples is approximately thesame, it can be assumed that high-temperatureexposure by either annealing or ECAE processingdid not result in increased oxygen level within thesamples. As special precautions were taken to mini-mize or prevent oxidation during milling, as well asthe handling of the powders, it can be assumed thatthe bulk oxygen content within the samples musthave been incorporated from inherent oxygen withinthe bulk or on the surface of the starting powders. Itis expected that, during high-energy ball milling thisnative oxygen was forced into solution within the Cumatrix along with the dissolved Ta. Therefore, thenext level of complexity to address is whether theoxygen is uniformly distributed throughout the bulkalloy or preferentially segregates to particular fea-tures in the microstructure.

The first noteworthy detail observed for all thesamples was that oxygen was present only as a com-plex ion, such as TaO and CuO, and not as elementaloxygen. The substantially higher presence of tanta-lum oxides compared to CuO was anticipated, con-sidering the high affinity Ta and O have for eachother in comparison toCu andO. Specifically, theCu-1Ta alloy had an average cluster composition of30Cu-36Ta-34O with Ta ranging from 11 at.% to89 at.% and O ranging from 2 at.% to 53 at.%. The

statistics were tabulated using 85 clusters. In con-trast, the Cu-10Ta had an average composition of27Cu-50Ta-23O with Ta ranging from 30 at.% to93 at.% and O ranging from 0 at.% to 45 at.%. Thesecompositions seemvery reasonablewhen consideringthat the Ta content increases from 1% to 10%, but theoxygen content remained constant between the twoalloys. Since in these two cases the mass can be con-sidered conserved, the overall level of oxygen incor-porated into the clusters is expected to be lower inhigher Ta-containing alloys where the number den-sity of particles is higher. Looking at the two samplesof the Cu-10Ta alloy, those being as-milled andannealed at 723 K (450!C) for 1 h in forming gas, thecluster compositions were found to be 35Cu-57Ta-8O at.% and 29Cu-62Ta-9O at.%, respectively. TheTa ranges were from 44 at.% to 96 at.% and 46 at.%to 93 at.%, and the O ranges were 2 at.% to 18 at.%and 1 at.% to 20 at.%. Obviously, the Ta content wasclearly higher, and the oxygen content was signifi-cantly lower, for these conditions relative to the hightemperature ECAE condition. This finding indicatesthat oxygen contamination is more significant athigher temperatures, which can be explained as fol-lows. In the as-milled state, the oxygen concentrationwithin the clusters is low, whereas after exposure tothe higher temperatures of ECAE processing, thisoxygen content increases. This was shown to be thecase regardless of the bulk Ta concentration. Con-sidering the cryogenicnature of themillingprocess, itis unlikely that oxygen can diffuse to or be furtherincorporated into theTa clusters. As the temperatureis increased to 723 K (450!C), oxygen began to diffusethrough the Cu lattice. The Cu oxide can be reducedin the presence of the forming gas, but this is not truefor Ta oxide. Furthermore, the temperature of 723 K(450!C) corresponds to an extremely low homologoustemperature for Ta, "0.22TM. It is, therefore, expec-ted that Ta may not readily react with the availableoxygen or move through the lattice until highertemperatures are reached. Since in both of the ECAEprocessed samples, higher levels of oxygen werefound to be incorporated in the clusters, it expectedthat the temperature of 973 K (700!C) provides suf-ficient activation energy allowing for oxygen to dif-fuse and interactwith theTa clusters. Indeed, a smallexotherm has been observed near this temperaturerange in previously reported differential scanningcalorimetry (DSC) data.25 This interaction betweenTaandoxygen is likely facilitatedby the fast diffusionpathways provided by the existing GB networks.

Upon careful inspection of the APT data, newinsights related to the partitioning of oxygen in theTa-rich clusters were revealed. Initial APT resultsseemed to indicate that oxygen was located at thesurface of the Ta-rich particles. This findingrequired careful consideration since evaporationfield differences between Ta and O can skew theevaporation process, introducing artifacts into theAPT reconstructions. However, it is reasonable toexpect that such an evolutionary process could occur


2808

with significant consequences on the stability andcoherency of the resultant particles. Additional APTstudies, coupled with aberration-corrected STEMand HRTEM, are required to confirm this hypothe-sis. These studies are currently underway.

CONCLUSION

The TEM and APT results indicated that the size ofTa clusters was compositionally independent of thebulk Ta content of the samples; instead, the increasedTa content manifested itself in a higher number den-sity of Ta clusters. Bulk oxygen content was found tobe roughly the same and consistent between the twocompositions for the given processing method. Inter-estingly, the as-milled and 723 K (450!C) annealedsamples had considerable lower oxygen content in theclusters compared to those ECAE processed at 973 K(700!C). The higher temperature allowed the oxygento diffuse out of the solid solution from the Cu matrixand attach to the Ta clusters. This indicated thatoxygen contamination plays a more significant role athigher processing/annealing temperatures. In initialAPT observations, oxygen appeared to be located atthe surface of the Ta particles. This indicatedincreased oxygen incorporation into the clusters,occurring after their as-milled formation. These clus-ters also showed similar structures in the TEM, but amore in-depth study is required before a definitiveconclusion can be drawn.

The MD simulations confirmed the role played byTa clusters in preventing or at least minimizing thegrain growth. The simulations also confirmed theAPT results for the Cu-1Ta alloy26 with respect to theTa particle coarsening along GBs via an Ostwaldripening process. The particles remained essentiallythe same size, allowing them to remain effectivepinning sites for preventing grain growth (in additionto being obstacles to dislocation motion). It has beenhypothesized that the particles within the interior ofthe grains have only a limited amount of coarsening.The kinetics and limits of such coarsening are relatedto the bulk Ta concentration and relative position ofthe cluster within the microstructure.

ACKNOWLEDGEMENTS

M. Rajagopalan and K. N. Solanki are grateful forthe financial support for this work from the ArmyResearch Laboratory award number W911NF-15-2-0038 and would also like to thank the LeRoy EyringCenter for Solid State Science at Arizona StateUniversity. G. P. Purja Pun and Y. Mishin weresupported by the U.S. Army Research Office undercontract number W911NF-15-1-007.

REFERENCES

1. A.J. Detor and C.A. Schuh, Acta Mater. 55, 371 (2007).2. A.J. Detor and C.A. Schuh, Acta Mater. 55, 4221 (2007).3. A.J. Detor, J.K. Miller, and C.A. Schuh, Philos. Mag. 86,

4459 (2006).4. K.A. Darling, R.N. Chan, P.Z. Wong, J.E. Semones, R.O.

Scattergood, and C.C. Koch, Scripta Mater. 59, 530 (2008).

5. K.A. Darling, B.K. VanLeeuwen, C.C.Koch, R.O. Scattergood,Mater. Sci. Eng. A 527, 357 (2010).

6. K.A. Darling, B.K. VanLeeuwen, J.E. Semones, C.C. Koch,R.O. Scattergood, L.J. Kecskes, and S.N. Mathaudhu, Ma-ter. Sci. Eng. A 528, 4365 (2011).

7. J.M. Dake and C.E. Krill III, Scripta Mater. 66, 390 (2012).8. R.J. Perez, H.G. Jiang, C.P. Dogan, and E.J. Lavernia, Me-

tall. Mater. Trans. A 29, 2469 (1998).9. L. Shaw, H. Luo, J. Villegas, and D. Miracle, Acta Mater. 51,

2647 (2003).10. A.M. El-Sherik, D. Boylan, U. Erb, G. Palumbo, and K.T.

Aust, Mater. Trans. A 238, 727 (1992).11. K. Boylan, D. Ostrander, U. Erb, G. Palumbo, and K.T.

Aust, Scripta Metall. Mater. 25, 2711 (1991).12. A. Michels, C.E. Krill, H. Eharhardt, R. Birringer, and D.T.

Wu, Acta Mater. 47, 2143 (1999).13. P. Knauth, A. Charai, and P. Gas, Scripta Metall. Mater. 28,

325 (1993).14. C.C. Koch, R.O. Scattergood, K.A. Darling, and J.E. Se-

mones, J. Mater. Sci. 43, 7264 (2008).15. K.A. Darling, M.A. Tschopp, B.K. VanLeeuwen, M.A.

Atwater, and Z.K. Liu, Comp. Mater. Sci. 84, 255 (2014).16. T. Chookajorn, H.A. Murdoch, and C.A. Schuh, Science 337,

951 (2012).17. M. Saber, H. Kotan, C.C. Koch, and R.O. Scattergood, J.

Appl. Phys. 113, 063515 (2013).18. Z. Chen, F. Liu, H.F. Wang, W. Yang, G.C. Yang, and Y.H.

Zhou, Acta Mater. 57, 1466 (2009).19. J. Li, J. Wang, and G. Yang, Scripta Mater. 60, 945 (2009).20. H.A. Murdoch and C.A. Schuh, Acta Mater. 61, 2121 (2013).21. C.C. Koch, R.O. Scattergood, M. Saber, and H. Kotan, J.

Mater. Res. 28, 1785 (2013).22. K.A. Darling, E.L. Huskins, B.E. Schuster, Q. Wei, and L.J.

Kecskes, Mat. Sci. Eng. A. 638, 322 (2015).23. T. Frolov, K.A. Darling, L.J. Kecskes, and Y. Mishin, Acta

Mater. 60, 2158 (2012).24. K.A. Darling, M.A. Tschopp, R.K. Guduru, W.H. Yin, Q.

Wei, and L.J. Kecskes, Acta Mater. 76, 168 (2014).25. K.A. Darling, A.J. Roberts, Y. Mishin, S.N. Mathaudhu, and

L.J. Kecskes, J. Alloys Compd. 573, 142 (2013).26. T. Rojhirunsakool, K.A. Darling, M.A. Tschopp, G.P. Purja

Pun, Y. Mishin, R. Banerjee, and L.J. Kecskes. MRS Com-mun. 5, 333 (2015).

27. C.M. Muller, S. Parviainen, F. Djurabekova, K. Nordlund,and R. Spolenak, Acta Mater. 82, 51 (2015).

28. C.M. Muller, A.S. Sologubenko, S.S.A. Gerstl, and R.Spolenak, Acta Mater. 89, 181 (2015).

29. G.P. Purja Pun, K.A. Darling, L.J. Kecskes, and Y. Mishin,Acta Mater. 100, 377 (2015).

30. V.M. Segal, Mater. Sci. Eng. A 197, 157 (1995).31. M. Furukawa, Z. Horita, M. Nemoto, T.G. Langdon, and J.

Mater, Sci. 36, 2835 (2001).32. V.M. Segal, Mater. Sci. Eng. A 271, 322 (1999).33. R.Z.Valiev andT.G.Langdon,Prog.MaterSci.51, 881 (2006).34. Y.T. Zhu and T.C. Lowe, Mater. Sci. Eng. A 291, 46

(2000).35. M.K. Miller, Atom Probe Tomography (New York: Kluwer

Academic/Plenum Publishers, 2000). ISBN-13: 9780306464157; ISBN-10: 0306464152.

36. M.K. Miller and E.A. Kenik, Microsc. Microanal. 10, 336(2004).

37. R.P. Kolli and D.N. Seidman, Microsc. Microanal. 13, 272(2007).

38. F. Vurpillot, A. Bostel, and D. Blavette, Appl. Phys. Lett. 76,3127 (2000).

39. F. Vurpillot, A. Cerezo, D. Blavette, and D.J. Larson, Mi-crosc. Microanal. 10, 384 (2004).

40. B. Gault, F. De Geuser, L. Bourgeois, B.M. Gabble, S.P.Ringer, and B.C. Muddle, Ultramicroscopy 111, 683 (2011).

41. X. Sauvage, L. Renaud, B. Deconihout, D. Blavette, D.H.Ping, and K. Hono, Acta Mater. 49, 389 (2001).

42. A. Stukowski, Model Simul. Mater. Sci. Eng. 18, 015012(2010).


effect of ta solute concentration on the microstructural evolution...

Documents