orientation dependent recovery and recrystallization behavior of hot-rolled molybdenum

Int. Journal of Refractory Metals and Hard Materials 48 (2015) 179–186

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Int. Journal of Refractory Metals and Hard Materials

j ourna l homepage: www.e lsev ie r .com/ locate / IJRMHM

Orientation dependent recovery and recrystallization behavior ofhot-rolled molybdenum

S. Primig a,⁎, H. Clemens a, W. Knabl b, A. Lorich b, R. Stickler c

a Department of Physical Metallurgy and Materials Testing, Montanuniversität Leoben, Franz-Josef Straße 18, 8700 Leoben, Austriab PLANSEE SE, Metallwerk-Plansee-Straße 71, 6600 Reutte, Austriac University of Vienna, Department of Physical Chemistry, Währinger Strasse 42, 1090 Vienna, Austria

⁎ Corresponding author.E-mail addresses: [email protected] (S. Pr

[email protected] (H. Clemens), [email protected] (A. Lorich), roland.stickler@

http://dx.doi.org/10.1016/j.ijrmhm.2014.09.0080263-4368/© 2014 Elsevier Ltd. All rights reserved.

a b s t r a c t
a r t i c l e i n f o
Article history:Received 17 June 2014Accepted 5 September 2014Available online 16 September 2014

Keywords:MolybdenumHot-rollingAnnealingRecoveryRecrystallizationEBSD

Due to its outstanding physical and chemical properties, technically pure molybdenum is nowadays frequentlyused in electronics and coating-technology besides its traditional applications as high temperature material. Atpresent, the production of large plates which are used as sputtering targets is of great importance. These platesare processed by hot-rolling of sintered pre-material with intermediate recrystallization annealing treatments.In the present investigation we studied the microstructural and textural evolution during dynamic recovery andstatic recrystallization of molybdenum by using an industrially processed molybdenum plate which has beenhot-rolled to a degree of deformation N60% followed by static recrystallization annealing between 1000 and1300 °C for various holding times up to 10 h. For the microstructural and textural characterization electronchanneling contrast imaging and electron back-scatter diffraction were used. The orientation dependentsubgrain size, the stored energy during hot-rolling, and the growth of individual subgrains upon static annealingindicated that theα-fiber subgrains, which are the largest ones after hot-rolling, exhibit slower static coarseningkinetics than the γ-fiber grains which are initially much smaller but contain a higher amount of stored energy.Furthermore, subgrains of other orientations located at in-grain shearbands seem to have growth advantageswhich cause a weakening of the rolling texture. A sluggish late-stage recrystallization behavior due to the recov-ery controlled annealing behavior has been revealed.

© 2014 Elsevier Ltd. All rights reserved.

1. Introduction

Molybdenum is a refractory metal with a body-centered-cubic(bcc) crystal lattice, a high stacking fault energy stacking fault energyof 0.3 J/m2 and a melting point of TM = 2620 °C [1–3]. Besides its highmelting point, molybdenum exhibits some further outstanding physicaland chemical properties as for example a high thermal conductivity anda low thermal expansion [1]. Therefore, new fields besides the tradition-al high temperature applications have been established during the lastfew decades. Nowadays, technically pure molybdenum is mainly ap-plied at ambient temperatures as a high performance functional materi-al in electronics and coating technology [1,4]. Due to the increasing sizeof dimensions to be coated, especially large plates for sputtering targetsare required. These plates are processed by hot-rolling of sintered pre-material [4] with intermediate recrystallization annealing treatments.Homogenously fine-grained, defect-free plates with a uniform textureare demanded by the customers. However, especially the thermo-mechanical processing of the required large dimensions remains

imig),[email protected] (W. Knabl),univie.ac.at (R. Stickler).

challenging. Molybdenum tends to intergranular fracture during hot-deformation if the deformation parameters are not chosen carefully[5]. Especially an intermediate recrystallization annealing after the ini-tial hot-rolling passes seems to be crucial for the subsequent hot-formability.

Due to the high stacking fault energy of molybdenum [3] theelevated-temperature behavior is predominantly governed by recoveryprocesses. The nucleation of subsequent static recrystallization occursvia coalescence and growth of individual subgrains which have beenformed during dynamic recovery [6–9]. These are the only mechanismsto control the microstructural evolution during processing of molybde-num that does not undergo phase transformations in the solid state.Both recovery and recrystallization are strongly orientation dependentin this body-centered-cubic material which develops the α-fiber withamaximum at the rotated cube and γ-fiber textural components duringhot-rolling. This behavior is characteristic for rolling bccmetals [10–14].The α-fiber subgrains have been reported to have a low stored energyand require less than five slip systems for their deformation which ac-counts for their blurry “mosaic block structure” [15] appearance. In con-trast, the γ-fiber subgrains exhibit a high stored energy, require five slipsystems and exhibit a more distinct subgrain structure [16–18].

Previous microstructural studies on the hot-deformation and an-nealing behavior of sintered starting materials, which were deformed

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in uniaxial compression, have revealed that the evolution of the defor-mation texture is strongly dependent on the hot-working temperature[7–9]. With an increasing deformation temperature a shift from theb111N parallel to the loading direction (//LD) fiber to the b001N//LDfiber has been observed. This has been explained in terms of the increas-ing amount of dynamic recrystallization phenomena, such as continu-ous dynamic recrystallization or strain-induced boundary migration oflow-stored energy b001N//LD subgrains [9]. The static recrystallizationbehavior is determined by the orientation dependent growth character-istics of individual subgrains [7]. Due to the low amount of storedenergy available after dynamic and static recovery, the late-stage re-crystallization behavior is sluggish. This results in residual islandsof recovered b001N//LD subgrains even after long annealing times[9,19–21].

This investigation focuses on the static recrystallization behavior ofan industrially processed hot-rolled molybdenum plate between 1000and 1300 °C for various holding times up to 10 h. Electron channelingcontrast imaging (ECCI) [22,23] was applied for the microstructuralcharacterization regarding the kinetics of recrystallization as well asthe recrystallized grain size and shape. The subgrain structure of theas-rolled startingmaterial aswell as the orientationdependent subgraincoarsening behavior was studied by electron back scatter diffraction(EBSD) [24–26]. The results are discussed in comparison to the micro-structural and textural evolution during deformation and annealing ofhigh stacking fault energy bcc metals in general and to the resultsobtained during uniaxial compression of similar starting materials inparticular [7–9].

2. Materials and methods

The startingmaterial for the present investigationwas an industrial-ly processed hot-rolled plate of the current quality of technically puremolybdenum. The chemical composition as determined by glow dis-charge mass spectroscopy (GDMS) of this plate in its as-rolled state isgiven in Table 1. The same plate (denoted as “Sheet 1”) was used inthe grain boundary segregation studies of Babinsky et al. [27,28]. Thisplate was produced by cold-isostatic pressing of molybdenum powderfollowed by a sintering process above 1800 °C. The subsequent hot-rolling was carried out above 1000 °C to a degree of deformationN60% in order to store sufficient energy for a complete recrystallizationduring static annealing within technically relevant times [19]. The aver-age density of the plate was determined as 10.21 g/cm3which indicatesthat a small amount of residual porosity of approximately 0.7% remains(theoretical density of molybdenum: 10.28 g/cm3 [1]).

Specimens were cut from the center of the plate with dimensions ofapproximately 1 × 0.5 × 5 cm3. These specimens were recrystallizationannealed in a hydrogen atmosphere between 1000 and 1300 °C for

Table 1Chemical composition of the molybdenum plate inits as-deformed condition as determined by GDMS.“b” indicates a value below the detection limit.

Element [μg/g]

C b8.5O ≤12N b1.3K 4Ba 1.0P 2.4La 0.02Fe 3.7Si 1.1Mg 0.9Al 0.2Ca 0.2S 0.09W 120

various holding times between 1 min and 10 h followed by furnacecooling. The total heating time was 20 min. Vickers hardness measure-ments (HV 10) were carried out on the annealed specimens and thestarting material. Metallographic sections were prepared by standardgrinding and polishing followed by electropolishing according to theprocedure described in [29]. ECCI investigations were carried out in aZeiss Evo 50 scanning electronmicroscope using an acceleration voltageof 15 kV atworking distances of 10mmor below. The recrystallized vol-ume fractionwas determined using a point countingmethod. The linearinterceptmethod parallel to the rolling directionwas used to determinethe recrystallized grain size in rolling direction. The starting materialand several specimens annealed at 1100 °Cwere additionally examinedby EBSD with a FEI Versa 3D dual beam microscope equipped with anEdax Hikari XP EBSD camera. An area of 1.2 × 1.2 mm2 was scannedusing a hexagonal grid and a step size of 0.65 μm. An acceleration volt-age of 20 kV, a working distance of 18 mm and a 6 × 6 binning of theEBSD camera were used. The data evaluation was carried out with theTSL OIM Analysis 7 software. As data clean-up procedures, the confi-dence index standardization followed by a single grain dilation (toler-ance 5.0°, minimum size 10 data points) iteration were used. For thefurther data evaluation, only points with a confidence index of N0.1were considered. The orientation distribution function (ODF) cross-section at φ2 = 45°, which includes the location of the α-fiber and theγ-fiber component in rolled bcc materials [14], was calculated for thestarting material. Furthermore, the local average misorientation (LAM)map was determined for a reduced area of the starting material'sdataset. The 2nd next neighbor datapoints and a maximummisorienta-tion of 5° were used. For the evaluation of the orientation dependentsubgrain coarsening and recrystallization kinetics, the datasets of thestartingmaterial and the annealed conditions were divided into subsetswhich contain (i) all data, (ii) the α-fiber datapoints, (iii) the γ-fiberdatapoints, or (iv) all points which do not belong to one of these twomain fibers (denoted as “others” in the following). The tolerance anglefor this procedurewas 20°. The fraction, average grain size, and local av-erage misorientations were calculated for all the individual subsets. Agrain tolerance angle of 0.5° was used to calculate the (sub)grain sizes.It must be noted that the subgrain sizes determined by this procedureshould not be taken as absolute values due to misindexed points closeto the subgrain boundaries which are discarded by the data clean-up.The determination of the subgrain size by ECCI or by transmission elec-tronmicroscopy would yield different values due to differences in reso-lution. However, the orientation dependent subgrain characteristics canbe used for revealing general trends regarding the orientation depen-dent coarsening behavior. Furthermore, no size criterion was appliedto distinguish between grains and subgrains.

3. Results

Fig. 1 is an ECC image of the microstructure in the center of the as-rolled starting material, a hot-rolled molybdenum plate. RD and NDstand for the rolling and the normal direction, respectively. This imageshows a recovered subgrain structure with some residual micro-pores from the sintering process which may have widened duringelectropolishing. Even if it is sometimes difficult to clearly reveal high-angle grain boundaries with this technique, individual regions with dif-ferent appearances of their subgrain structure are discernible. The re-gion which is surrounded by the white line exhibits a rather blurryappearance while the region surrounded by the black line exhibits afar more distinct and equiaxed subgrain structure. Such areas are char-acteristic for the microstructure of hot-rolled molybdenum.

During static recrystallization annealing of specimens from the cen-ter of the same plate, individual subgrains grow or coalesce with theirimmediate neighbors leaving behind large, defect-free grains. Fig. 2shows the microstructural evolution during static annealing at1200 °C. Fig. 2a shows the microstructure of a specimen which wasannealed for 5 min. The major area of this micrograph still exhibits a

Fig. 1. Microstructure in the center of the as-rolled starting material, a molybdenumplate, studied by ECCI. RD and ND stand for the rolling and the normal direction, respec-tively. Residual micropores are indicated by arrows. Thewhite and the black lines indicateregions/grains with two different subgrain structures.

Fig. 2.Microstructural evolution of a hot-rolledmolybdenumplate during static annealingat 1200 °C studied by ECCI. (a) Was annealed for 5 min. Two small and one larger recrys-tallized grains are marked by the arrows. (b) Was annealed for 1 h and (c) for 10 h. In(b) and (c) some small non-recrystallized islands are marked by arrows.

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recovered subgrain structure. However, some small and several larger(marked by the arrows) recrystallized grains are visible. These recrys-tallized grains sometimes exhibit quite irregular shapes. The specimenin Fig. 2b is more or less fully recrystallized after 1 h of annealing. Thegrain boundaries are slightly serrated. Some small non-recrystallizedislands which are marked by the arrows are still present. Such areasare even present after 10 h annealing (Fig. 2c). The grains in this micro-graph are slightly larger. The grain boundary serrations are still present,however, they appear less distinct as in Fig. 2b. Residual micropores canbe found in all specimens.

The kinetics of recrystallizationwas studied by evaluating numerousof such ECCImicrographs. The results aswell as the corresponding hard-ness values are shown in Fig. 3. Fig. 3a shows the recrystallized volumefraction as determined by a point counting method. This graph revealsthat at 1000 °C annealing time, almost no recrystallized grains arepresent even after 2 h of annealing. At 1300 °C the entire microstruc-ture, except some small non-recrystallized islands, is recrystallizedeven after the shortest annealing times. In between these two temper-atures recrystallization occurs after a certain period of incubation, how-ever, the recrystallization behavior becomes more and more sluggishafter approximately 80% recrystallization. The mean hardness values,Fig. 3b, decrease from 216 HV 10 in case of the shortest annealingtimes at the lowest temperatures to 165 HV 10 after 2 h annealing at1300 °C. The recrystallized grain size (Fig. 3c) ranges from 50 μm//RDwhen only a small fraction of the microstructure is recrystallized to170 μm in all conditions where the microstructure is more or less fullyrecrystallized. There is some grain growth after recrystallization in thecase of the long-time annealing for 10 h at 1200 °C.

The starting material and several specimens annealed at 1100 °Cwere additionally examined by EBSD in order to reveal the orientationdependent subgrain coarsening and subsequent recrystallization kinet-ics. Fig. 4 shows the EBSD results obtained in the center region of thehot-rolled startingmaterial. Fig. 4a is the inverse pole figure (IPF) color-ingmapwith crystal directions in the normal direction of the plate. Thehigh-angle boundarieswith amisorientation angle of N15° are shown asblack lines. This image reveals that there are many grains with either ab001N direction parallel to the normal direction (//ND) or b111N//ND. Furthermore,many in-grain shearbandswith high localmisorienta-tions inside the original grains can be found. Even some additional seg-ments of high-angle boundaries have been formed during rolling. Fig. 4bis the φ2 = 45° cross-section of the orientation distribution function(ODF) calculated using the scan data shown in Fig. 4a. It evidences the

presence of the characteristic α-fiber with a maximum at the rotatedcube and the γ-fiber components which are formed during hot-rollingof bcc metals. Fig. 4c is the IPF coloring map of the upper left corner ofthemap shown in Fig. 4a without drawing the high-angle grain bound-aries. This image reveals differences in the appearance of the individualgrains of each textural component. Some b001N//ND grains rather ex-hibit a blurry appearance which indicates small continuous rotation of

Fig. 3. Recrystallization behavior during static annealing of a hot-rolled molybdenumplate. (a) Recrystallized volume fraction, (b) hardness (HV 10) and (c) recrystallizedgrain size versus annealing time.


the local orientation while other regions consist of coarse subgrains. Incontrast, the b111N//ND grains in general have a much more distinctstructure offine recovered subgrains. Three arrows indicate the positionsof the point to point misorientation plots shown in Fig. 4e. Fig. 4d is theLAM map of the area shown in Fig. 4c. It shows qualitatively the localstrain which is especially high in regions exhibiting in-grain shearbands.Fig. 4e is a graph of the point to point misorientation plots of the threepositions marked in Fig. 4c. It points out the orientation dependent dif-ferences in the local misorientation dependent on the average orienta-tion of the grains. The curve which was taken inside a b001N//NDgrain showsonlyminor variations in the localmisorientation. In contrast,the curve taken in a b111N//ND region exhibits some small peaks andtwo high-angle grain boundaries. The curve in the region of theshearbands (SB) exhibits multiple peaks of high local misorientations.

Fig. 5 shows the EBSD IPF coloring maps in the center regions of thespecimens annealed at 1100 °C for 5 min (Fig. 5a), 15 min (Fig. 5b) and2 h (Fig. 5c). In Fig. 5a there are no recrystallized grains or small coa-lesced regions discernible by the naked eye. However, some subgraincoarsening has already taken place which will be shown in the follow-ing. Fig. 5b exhibits many recrystallized grains without a recoveredsubgrain structure. It must be noted that these recrystallized grainshave more random orientations than the grains shown in Figs. 4a and5a. Themicrostructure in Fig. 5c is close to fully recrystallized. However,it is obvious, that there are still some small non-recrystallized islandsleft as already discernible in Fig. 2c. These islands always exhibit orien-tations close to b001N//ND.

The datasets of Figs. 4a and 5were taken to calculate the orientationdependent volume fraction, average (sub)grain size and average localmisorientation values shown in Fig. 6.

Fig. 6a shows the volume fraction of each textural component as afunction of the annealing time. Similar to Fig. 4b, this graph indicates ahigh fraction of α-fiber in the starting material. During annealing thefraction of α-fiber decreases from 57% in the starting material to 37%after 2 h at 1100 °C. In contrast, the fraction of the “others”, i.e. all re-gions which do not belong to the two main fiber textures, increasesfrom 28% in the starting material to 54% during annealing. The fractionof the γ-fiber increases during the initial 15 min from 8% to 11%. How-ever, after annealing for 2 h, the fraction decreases again to 7%. Fig. 6billustrates the orientation dependent coarsening behavior. As visible inFig. 5a, there is just subgrain coarsening, but no obvious formation oflarger, defect-free grains after 5 min of annealing at 1100 °C. After15 min of annealing, the major area of the EBSD scan is already recrys-tallized, however, there are also some areas of recovered subgrainspresent. In this case the grain size criterion does not distinguish be-tween grains and subgrains but determines the average grain size ofboth kinds of crystallites. This graph does not contain the data of theclose to fully recrystallized specimen shown in Fig. 5c. Due to the largerecrystallized grain size, the data of this scan is not significant anymore.Fig. 6b reveals that initially the α-fiber subgrains are the largestones (mean value: 8.4 μm2), followed by the others (7.2 μm2) and theγ-fiber subgrains (6.12 μm2). After 5 min of annealing, there is somesubgrain coarsening of all textural components. However, after 15 minannealing time, the orientation dependent (sub)grain size sequence re-verses. Theα-fiber (sub)grains are the smallest ones (14.8 μm2) follow-ed by the others (24.8 μm2) and the γ-fiber (sub)grains (31.7 μm2).Fig. 6c shows the orientation dependent LAM values. These are initiallylower in the α-fiber (mean value: 0.9°) than in the others (0.93°)followed by the γ-fiber (0.95°). After 15 min of annealing, the LAMvalues decrease due to recovery and recrystallization processes andthe sequence is again reversed. The values in theα-fiber are the highestones (0.69°) followed by the others (0.57°) and by the γ-fiber (0.52°).After 2 h of annealing, the sequence is equal but the values have de-creased further. Similar results with different kinetics have been obtain-ed at all annealing temperatures examined.

4. Discussion

It is well known that the static recrystallization behavior of metalsstrongly depends on the as-deformed starting microstructure [30,31].The recrystallization behavior of molybdenum in particular is stronglydominated by concurrent recovery processes [3,6–9,32]. In this case itdepends on the orientation dependent subgrain characteristic [9,33]. Inthe following paragraphs the global as well as the orientation dependentstatic subgrain coarsening and recrystallization kinetics of an industriallyprocessed hot-rolled molybdenum plate will be discussed. The findingswill be compared to bccmetals in general and to previous results obtain-ed during static annealing subsequent to uniaxial compression with acomparable degree of deformation at different temperatures [7–9].

The overall recrystallization kineticswas studied byusing ECCI [22,23]micrographs as shown in Figs. 1 and 2. Fig. 1 is the microstructure of the

Fig. 4. EBSD results of thehot-rolled startingmaterial. (a) IPF coloringmapwith crystal directions in the normal direction of the plate (see inset). All boundarieswith amisorientation angleof N15° are shown as black lines. (b) φ2 = 45° section of the ODF calculated using the scan data shown in (a). (c) IPF coloringmap of the upper left corner of the map shown in (a). Threearrows indicate the positions of the point to pointmisorientation plots shown in (e). (d) LAMmap of the area shown in (c)with high-angle grain boundaries shown as black lines. Refer tothe inset for the corresponding color code. (e) Point to point misorientation plots of the three positions marked in (c).


hot-rolled starting material indicating regions/grains with differentsubgrain characteristics. Some regions as indicated by the white line ex-hibit a blurry appearance similar to deformation bands (“mosaic blockstructures”) [15,34], while other regions (black line) exhibit a more dis-tinct subgrain structure [16–18]. It is supposed that initially more strainis accommodated by the former grains/regions [15,35]. Upon staticannealing of this starting material (Fig. 2) there is no nucleation of

new crystallites but growth and/or coalescence of individual subgrains[36,37]. However, only very few of these subgrains have such growth ad-vantages [33,38] that they can become new defect-free recrystallizedgrains (as marked by the arrows in Fig. 2a) [30,39]. Furthermore, suchgrains exhibit rather unusual shapes with serrated grain boundaries[40,41] compared to, e.g., recrystallized austenitic microstructures [31].In Fig. 2b a more or less fully recrystallized microstructure with serrated

Fig. 5. (a) IPF coloringmapwith crystal directions in the normal direction (see inset of Fig. 4a) of the specimenwhichwas annealed at 1100 °C for 5min, (b) 15min and (c) for 2 h. For theIPF coloring code refer to the inset of Fig. 4a. All boundaries with a misorientation angle N 15° are shown as black lines.


grain boundaries and some small non-recrystallized islands is shown.After 10 h of annealing at 1200 °C (Fig. 2c) these islands are still presentwhile the grain boundaries appear slightly less serrated. It is assumed thatthe serrated grain boundaries in Fig. 2b are the results of the progressionof the individual recrystallization fronts by subgrain coalescence andgrowth. Only after complete recrystallization, there is time for straighten-ing of the grain boundaries driven by the minimization of the interfacialenergy [38,42]. The recrystallization kinetic studies and the hardnessmeasurements (Fig. 3a and b) indicate more or less complete recrystalli-zation at 1300 °C and almost no recrystallization at 1000 °C in case of allannealing times studied. The annealing temperatures in between 1000and 1300 °C indicate Johnson–Mehl–Avrami–Kolmogorov- (JMAK-)[43] like recrystallization kinetics after a certain period of incubation.However, the late-stage recrystallization behavior is sluggish and cannotbe described by a standard JMAK-approach [20,21]. The recrystallizedgrain size of themore or less fully recrystallizedmicrostructure is approx-imately equal in all conditions except the condition annealed for 10 h andis independent on the annealing temperature (Fig. 3c) [31,44]. An activa-tion energy of 4.4 eV for primary recrystallization of a similar startingma-terial has been determined by Hünsche et al. in [45].

A more detailed discussion of the recrystallization in hot-rolled mo-lybdenum can be achieved by examining the EBSD data. Fig. 4a and bshows that the starting material exhibits the characteristic texture

that is obtained during rolling of bcc metals with the α-fiber and itsmaximum at the rotated cube as well as the γ-fiber textural compo-nents [10–14]. By examining the IPF coloring map (Fig. 4a) the regionindicated by the white line in Fig. 1 with the blurry appearance are α-fiber grains while the regions with the more distinct subgrain structure(black line) belong to theγ-fiber. In a similarmanner, blurry b001N//LDand distinct b111N//LD grains have been found in molybdenum de-formed in uniaxial compression [8,9]. In both cases the former ones ex-hibit low Taylor factors and, thus, low stored energies, while the latterones exhibit high Taylor factors and, therefore, high stored energies[16–18,46]. Due to this fact these components as well as the “others”in rolled and compressively deformed material can be compared toeach other. The main difference between rolling and uniaxial compres-sive deformation is the far more distinct in-grain shearband formation[47,48] which appears frequently in “others” grains in the hot-rolledmaterial, possibly due to the absence of the rotational symmetry of uni-axial compression. The high local misorientations and the formation ofadditional segments of high-angle grain boundaries [9,44,47] in suchgrains are also evidenced by the LAM values which are shown inFig. 4d. Furthermore, the difference in local misorientation in all threetextural components is indicated by the point-to-point misorientationplots in Fig. 4e. The curve of the α-fiber b001N//ND grain has onlyvery small peaks but shows increasing values in the case of the γ-fiber

Fig. 6.Results of the orientation dependent EBSD data evaluation: (a) evolution of the vol-ume fraction of the individual textural components, (b) average grain size (an area takinginto account both grains and subgrains), and (c) LAM as a function of annealing time at1100 °C.


b111N//ND andmaximumvalues in case of the “others” grainswith thein-grain shear-bands. The orientation dependent subgrain characteris-tics in Fig. 6 show that similar to uniaxial compression and annealingat 1100 °C and for bcc metal in general [7] the α-fiber subgrains in thehot-rolled starting material are initially the largest ones followed bythe “others” and the γ-fiber subgrains. In contrast, the highest LAMvalues can be found in the γ-fiber followed by the “others” and the α-fiber. In the starting material, the fraction of the α-fiber is the largestfollowed by the others and the γ-fiber which is also indicated in Fig. 4b.

The IPF coloring map obtained after annealing for 5 min at 1100 °Cas shown in Fig. 5a has a similar appearance to the starting material ifexamined by the naked eye. However, the orientation dependent

subgrain characteristics evidence subgrain coarsening [9,33] prior tothe onset of static recrystallization of all textural components in a simi-larmanner after 5min of annealing. The average subgrain sizes increaseslightly while the LAM values decrease. This corresponds to the periodof incubation [30,36,49] as can be seen in Fig. 3a.

After 15min of annealing (Fig. 5b) partial recrystallization has takenplace. In this case the orientation dependent (sub)grain characteristicsare reversed. The γ-fiber crystallites are the largest ones followed bythe “others” and the α-fiber. This indicates that the onset of growth oftheγ-fiber crystallites occurs early due to the high amount of stored en-ergy available, while the growth of the initially largeα-fiber subgrains isslow [16–18,46]. This is also similar to uniaxial compression and anneal-ing and indicates that stored energy is more important than the initialsubgrain size. After 15min of annealing the sequence of the LAM valuesalso reverses. The α-fiber exhibits the highest LAM values followed bythe “others” and the γ-fiber. This correlates with the stored energy con-sumed for (sub)grain growth of the individual components. The fractionof theα-fiber decreases, while the one of the “others” increases and thefraction of the γ-fiber remains more or less constant despite the earlyonset of growth in this fiber. This indicates that nucleation in the γ-fiber preferably occurs inside γ-fiber grains. Furthermore, duringgrowth preferably adjacent as-deformed γ-fiber regions are consumed[36,39,50].

After 2 h of annealing at 1100 °C (Fig. 5c) andmore or less completerecrystallization the sequence in LAM values is lower but stillequal, however, the fraction (Fig. 5a) of the individual componentshas changed. The “others” now are the largest fraction followed by theα-fiber. The γ-fiber fraction has decreased slightly. This indicates thatduring the late stages of recrystallization, similar to uniaxial compres-sion and annealing, the grains that belong to the “others” rather thanthe γ-fiber seem to exhibit maximum growth. However, due to thepresence of distinct in-grain shearbandswith high local misorientations[36,47,48,51], this effect is far more pronounced here. Therefore, therecrystallization texture is less strong than the rolling texture. Further-more, it can be seen that similar to uniaxial compression and annealing,the non-recrystallized islands belong the α-fiber. This kind of sluggishlate-stage recrystallization behavior is due to the low amount of storedenergy of such orientations and, thus, causes the deviations from theJMAK-kinetics [20,21,43].

As a consequence, it can be stated that the static recrystallization be-havior of this hot-rolled molybdenum plate is in many ways similar tothe results obtained during uniaxial compression and annealing at1100 °C. However, the explicit temperature evolution during rolling ofthe sinter-plate is similar but not exactly equal to compression speci-mens. Obviously, the main difference is the presence of in-grain shear-bands after rolling which are preferred nucleation sites due to thehigh local misorientations and additional segments of mobile high-angle boundaries introduced during rolling [47,48]. Therefore, it is pos-sible to weaken the rolling texture due to growth advantages ofsubgrains which do not belong to the two main textural componentsduring annealing [9]. Furthermore, it is assumed that such “fresh”high-angle boundaries have a higher mobility than the original high-angle boundaries because they do not have to drag as many impurities(e.g. see Table 1). Recently, both grain boundaries with andwithout im-purities in the same material have been examined by atom probe to-mography [27]. Even if more detailed studies are still required, it couldbe assumed that the former ones are original boundaries and the latterones are boundaries introduced during deformation. However, the dragof impuritiesmight also be the reasonwhy the straightening of serratedgrain boundaries as well as further growth after primary recrystalliza-tion are also quite sluggish.

5. Summary

ECCI and EBSD have been applied for the overall and orientationdependent study of the recrystallization kinetics of a hot-rolled


molybdenum plate. From this study the following conclusions can bedrawn:

• After rolling, the plate exhibits a rolling texture and orientation de-pendent subgrain structure which are characteristic for bcc metals.

• Theα-fiber subgrains, which are the largest ones after hot-rolling, ex-hibit slower static coarsening kinetics upon annealing than theγ-fibersubgrains which are initially much smaller but contain a higheramount of stored energy.

• Subgrains located at in-grain shearbands seem to have growth advan-tages which cause a weakening of the rolling texture.

• A sluggish late-stage recrystallization behavior due to the recoverycontrolled annealing behavior has been revealed.

• The straightening of the high-angle grain boundaries after primary re-crystallization is sluggish, possibly due to amechanism invoked by thedrag of impurities.

References

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Dr. Sophie Primig: Is the head of the “High Performance Materials and Advanced Steels”group at the Department of Physical Metallurgy and Materials Testing at theMontanuniversität Leoben. She finished her Ph.D. thesis “Recovery and RecrystallizationBehavior of Technically Pure Molybdenum” which was carried out in collaboration withPlansee SE in 2012. Today her work is mainly focused on steels, nickel-base alloys, molyb-denumandmolybdenumalloys. She is interested inmaterials characterization by electronmicroscopy, EBSD, thermal analysis, and atom probe tomography.

Prof. Helmut Clemens: Is the head of the Department of Physical Metallurgy andMaterials Testing at the Montanuniversität Leoben. The main focus of his work is the de-velopment, characterization and testing of advanced high-performance materials. Afterhis Ph.D. at theMontanuniversität Leoben heworked at Plansee SE, where hewas respon-sible for the development of novel titanium-aluminide alloys. After several years inGermany he came back to Leoben in 2003 as professor for Physical Metallurgy. Besidesteaching, his research is still focused on titanium aluminides, refractorymetals, and steels.

Dr.WolframKnabl: Is the head of the R&D department at Plansee SE Business Unit Indus-tries, an Austrian company which has been processing refractory metals and their alloysfor more than 90 years.

Dr. Alexander Lorich: Is a staff engineer and project manager in the R&D department atPlansee SE, where he is working in the field of thermo-mechanical treatment of refractorymetals.

Prof. Roland Stickler: Is emeritus professor of the Department of Physical Chemistry atthe University of Vienna. He spent several years of his professional career at theWesting-house R&D laboratories in Pittsburgh, Pennsylvania before coming back to Austria in thenineteen-seventies. His research has been in close collaboration with Plansee SE eversince.

orientation dependent recovery and recrystallization behavior of hot-rolled molybdenum

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