Microstructural Evolution of Dy2o3-Tio2 Powder Mixtures During Ball Milling and Post-Milled Annealing

materials

Article Microstructural Evolution of Dy2O3-TiO2 Powder Mixtures during Ball Milling and Post-Milled Annealing

Jinhua Huang, Guang Ran *, Jianxin Lin, Qiang Shen, Penghui Lei, Xina Wang and Ning Li

College of Energy, Xiamen University, Xiamen 361102, China; [email protected] (J.H.); [email protected] (J.L.); [email protected] (Q.S.); [email protected] (P.L.); [email protected] (X.W.); [email protected] (N.L.) * Correspondence: [email protected]; Tel./Fax: +86-592-2185278

Academic Editor: Jan Ingo Flege Received: 30 October 2016; Accepted: 23 December 2016; Published: 28 December 2016

Abstract: The microstructural evolution of Dy2O3-TiO2 powder mixtures during ball milling and post-milled annealing was investigated using XRD, SEM, TEM, and DSC. At high ball-milling rotation speeds, the mixtures were ﬁned, homogenized, nanocrystallized, and later completely amorphized, and the transformation of Dy2O3 from the cubic to the monoclinic crystal structure was observed. The amorphous transformation resulted from monoclinic Dy2O3, not from cubic Dy2O3. However, at low ball-milling rotation speeds, the mixtures were only ﬁned and homogenized. An intermediate phase with a similar crystal structure to that of cubic Dy2TiO5 was detected in the amorphous mixtures annealed from 800 to 1000 ◦C, which was a metastable phase that transformed ◦ to orthorhombic Dy2TiO5 when the annealing temperature was above 1050 C. However, at the same annealing temperatures, pyrochlore Dy2Ti2O7 initially formed and subsequently reacted with the remaining Dy2O3 to form orthorhombic Dy2TiO5 in the homogenous mixtures. The evolutionary mechanism of powder mixtures during ball milling and subsequent annealing was analyzed.

Keywords: microstructure; ball milling; dysprosium oxide; neutron absorber; phase evolution

1. Introduction High-energy ball milling has been widely used to prepare various types of materials, such as supersaturated solid solutions, metastable crystalline materials [1], quasicrystal phases [2], nanostructured materials [3], and amorphous alloys [4]. The technology was initially used in place of blending and sintering at elevated temperatures to prepare ceramic-strengthened alloys [5,6]. A large amount of mechanical energy is transformed into intrinsic energy in the target materials, which induces the formation of numerous defects in the crystal structure, such as vacancies, interstitials, cavities and dislocations, which are always in a non-equilibrium state [7–9]. The defects and structural disorders will increase the mobility of atomic diffusion and induce chemical reactions amongst components that are not present under equilibrium conditions [10]. Therefore, based on its excellent characteristics, ball milling was used to prepare bulk Dy2TiO5, which can be used as a neutron absorber in control rods in nuclear power plants. Control rods are very important in both operating and accident conditions because the nucleon reactivity must be controlled in order to safely operate a nuclear reactor [11]. In fact, bulk Dy2TiO5 prepared by ball milling and sintering has been used in Russian power plant water reactors, such as MIR and VVER-1000 RCCAs [12,13], because of the excellent nucleon characteristics of the element dysprosium, as natural dysprosium consists of ﬁve stable isotopes with high thermal neutron absorption cross sections. The decay products are Ho and Er, which are also able to absorb neutrons. All of the radionuclides have low gamma activity and short half-life periods. The absorption

Materials 2017, 10, 19; doi:10.3390/ma10010019 www.mdpi.com/journal/materials Materials 2017, 10, 19 2 of 12 cross sections of dysprosium isotopes range from 130 barn to 2600 barn. The region of resonance absorption is 1.6–25 eV, in which the absorption cross-section can reach approximately 1000 barn [14]. According to its equilibrium phase diagram, Dy2TiO5 has three crystal structural types depending 1350 ◦C 1680 ◦C on the temperature, orthorhombic ↔ hexagonal ↔ cubic [15], which have different physical properties and radiation resistance abilities. In fact, bulk Dy2TiO5 in the cubic crystal structure has the lowest neutron irradiation swelling and highest irradiation resistance. Therefore, it is necessary to synthesize bulk Dy2TiO5 in the cubic crystal structure. Jung [16] synthesized bulk Dy2TiO5 with high purity and density using a polymer carrier chemical synthesis process, in which ethylene glycol was used as an organic carrier for metal cations. An amorphous phase was detected below 800 ◦C, ◦ and orthorhombic Dy2TiO5 was observed after sintering for 1 h at 1300 C, while little else was observed while sintering in the range of 800 to 1300 ◦C. Panneerselvam [17] used both solid-state synthesis and wet chemical synthesis to prepare Dy2TiO5. However, the effect of sintering temperature on the phase evolution needs further investigation. The sinterability of Dy2O3 and TiO2 with different molar ratios was determined for various ball-milling and sintering conditions [18]. Amit Sinha [19] reported the synthesis of bulk Dy2TiO5 from mixtures of equimolar Dy2O3 and TiO2 powders in a two-step process: (I) pyrochlore Dy2Ti2O7 was initially formed and (II) Dy2Ti2O7 then reacted with the remaining Dy2O3 to form orthorhombic Dy2TiO5. The powder mixtures used in sintering were simply mixed during ball milling. Garcia-Martinez [7] observed the experimental phenomena of the transformation of Dy2O3 from cubic to monoclinic and the synthesis of a hexagonal high-temperature phase, reported as Dy2TiO5, in an equimolar Dy2O3-TiO2 mixture during ball milling. Therefore, further investigation is needed into the evolutionary behavior of the microstructure under different ball-milling conditions and the effect of the state of the ball-milled powder on the sintering behavior. In the present work, the microstructural evolutionary behavior and corresponding reaction mechanism of Dy2O3-TiO2 powder mixtures under two types of ball-milling parameters were investigated. The annealing behavior of the ball-milled mixtures was also examined.

2. Experiments

Powders of Dy2O3 (cubic crystal structure) and TiO2 (rutile crystal structure) with an average particle diameter of 5 µm and 50 nm, respectively, were used as raw materials. The raw powders of Dy2O3 and TiO2 were purchased from Beijing HWRK Chem Co., Ltd. (Beijing, China). The purity of the raw powders of both Dy2O3 and TiO2 was 99.9%. Ball milling of the molar fraction Dy2O3-50% TiO2 (Dy2O3:TiO2 = 1:1) powder mixtures was carried out on an SFM-1 high-energy planetary ball mill at room temperature. Stainless steel balls that were 5 mm in diameter were used as the milling media. The ball-to-powder mass ratio was 10:1, and the rotational speed was 200 rpm and 500 rpm. No more than one weight percent stearic acid was added to the powder mixtures as a process control agent to prevent excessive cold welding and aggregation amongst the powder particles. During ball milling, a 5-min stopping interval was used after milling for 55 min to prevent excess heat generation, which has an obvious effect on the ball-milling procedure. The powder mixtures used for microstructural analysis were extracted from the loose powders in the steel can, not from powders adhered to the surface of the stainless balls or the steel can wall, after ball milling for 4, 12, 24, 48, and 96 h. After various milling times, a small amount of ball-milled powders taken from the container were characterized and analyzed by X-ray diffraction (XRD) on a Rigaku Ultima IV X-ray diffractometer (Rigaku, Tokyo, Japan) with Cu Kα radiation (λ = 0.1540598 nm) and transmission electron microscopy (TEM) on a JEM-2100 instrument (JEOL, Tokyo, Japan). Analysis was also carried out for ball-milled powder mixtures annealed at different temperatures. The grain size was calculated using Suryanarayana and Grant Norton’s formula [20].

Kλ B cos θ = + η sin θ (1) r L Materials 2017, 10, 19 3 of 12 where, K is a constant (with a value of 0.9); λ is the wavelength of the X-ray radiation; L and η are the grain size and internal strain, respectively; and θ is the Bragg angle. Br is the full width at half-maximum (FWHM) of the diffraction peak after instrumental correction and can be calculated from the following equation: B = Br + Bs (2) where, B and BS are the FWHM of the broadened Bragg peaks and the standard sample’s Bragg peaks, respectively. The ball-milled mixtures and annealed mixtures were ﬁrst put in ethyl alcohol, and then adequately dispersed by ultrasonic vibration. A carbon-coated copper grid was used to collect the dispersed powders in the ethyl alcohol and then dried by ultraviolet lamp. After that, the prepared samples were observed by TEM. Differential scanning calorimetry (DSC) was used to analyze the thermal behavior of ball-milled powders at a 5 ◦C/min heating rate in argon atmosphere using a SAT 449C instrument (NETZSCH, Bavarian State, Germany ). The powder mixtures milled for 96 h were annealed at temperatures ranging from 700 to 1150 ◦C in a tube furnace under atmospheric conditions. The heating and cooling rates were both 5 ◦C/min.

3. Results and Discussion

The XRD patterns of Dy2O3-TiO2 powder mixtures milled at 500 rpm and 200 rpm for different times are shown in Figure1. The XRD results show that the crystal structure of the original Dy 2O3 phase and TiO2 phase are cubic and rutile, respectively. At the condition of 500 rpm, the diffraction peaks of cubic Dy2O3 and TiO2 broadened significantly and reduced in intensity with increased milling time. The broadening of the X-ray diffraction peaks is associated with the refinement in grain size and lattice distortions. Meanwhile, the diffraction peaks of monoclinic Dy2O3 can be observed in the X-ray patterns as indicated by the black inverted triangles in Figure1a. Ball milling induces a Dy2O3 phase transformation from the cubic to the monoclinic crystal structure. A broad, singular diffraction peak is also present that indicates the formation of the amorphous phase during ball milling. Interestingly, the amorphous peak is present at the location of the diffraction peak for the monoclinic Dy2O3 phase, not at the location of the diffraction peak for the cubic Dy2O3 phase, which indicates that the formed amorphous phase is derived from the monoclinic Dy2O3 phase, not from the cubic Dy2O3 phase. The transformation from cubic to monoclinic increases with increased milling time. After milling for 96 h, only the amorphous phase can be observed, which indicates that the monoclinic Dy2O3 phase was fully converted to the amorphous phase. Additionally, this behavior indicates that no new compounds are synthesized during ball milling. Even if new compounds were formed in the milled powders, the amount is very low and does not reach the sensitivity range of the X-ray measurement. Therefore, in the present work, the evolution of Dy2O3-TiO2 powder mixtures is as follows: ball milling first induces the transformation of Dy2O3 from the cubic to the monoclinic crystal phase, then monoclinic Dy2O3 undergoes amorphization, and finally the powder mixtures completely transform to the amorphous phase. However, the ball-milling behavior of powder mixtures at 200 rpm is distinctly different from that at 500 rpm. The change of the diffraction peaks with increased milling time at 200 rpm is shown in Figure1b. Although the diffraction peaks of cubic Dy 2O3 and TiO2 are also broadened and reduced in intensity with increased milling time, the diffraction peaks of TiO2 can be observed in the XRD spectrums and are not disappeared. After ball milling for 96 h, the intensity of diffraction peaks of Dy2O3 and TiO2 are also high. The powder mixtures are not changed completely to amorphization. According to the shape of XRD diffraction spectrums, the effect of ball milling on powder mixtures after milling for 96 h at 200 rpm is only similar to that after milling for 4 h at 500 rpm. Therefore, at low ball-milling rotation speeds, the powder mixtures are only fined and homogenized. Our experimental results are different from the results of G. Garcia-Martinez [7]. In their research, ball milling induced a phase transformation in Dy2O3 from cubic to monoclinic. However, a Dy2TiO5 compound with a hexagonal crystal structure was formed simultaneously. The ball-milled powders Materials 2017, 10, 19 4 of 12

Materials 2017, 10, 19 4 of 13 consisted of mixed phases of hexagonal Dy2TiO5 and monoclinic Dy2O3. The Dy2O3-TiO2 powder mixtures did not completely transform to the amorphous phase, but instead produced the hexagonal powder mixtures did not completely transform to the amorphous phase, but instead produced the Dy2TiO5 phase. This difference can be attributed to the different ball-milling conditions used in this hexagonal Dy2TiO5 phase. This difference can be attributed to the different ball-milling conditions research,used in with this research, special attention with special to the attention different to ball-millingthe different facilities.ball-milling During facilities. ball During milling, ball experimental milling, parametersexperimental such parameters as rotation such speed, as ball-millingrotation speed, time, ball-milling ball-milling time, media, ball-milling and the media, ball-to-powder and the ball- mass ratioto-powder have an mass important ratio have inﬂuence an important on the ball-milled influence on products the ball-milled even when products using even the when same using proportion the andsame type proportion of oxides. and For type example, of oxides. Gajovi´creported For example, Gajovi thatć reported nanosized that ZrTiO nanosized4 formed ZrTiO in4 formed ZrO2-TiO in 2 powderZrO2-TiO mixtures,2 powder whereas mixtures, only whereas amorphous only amorphous mixtures were mixtures obtained were duringobtained ball during milling ball in milling Stubiˇcar’s in workStubi [21čar’s,22]. work In addition, [21,22]. theIn addition, polymorphic the polymorphic transformation transformation of Ln2O3 was of Ln also2O observed3 was also in observed Gd2O3-TiO in 2 andGd Y22OO3-TiO3-2TiO2 and2 powder Y2O3-2TiO systems2 powder [23]. systems [23].

Figure 1. X-ray diffraction patterns of the powder mixtures milled at (a) 500 rpm and (b) 200 rpm for Figure 1. X-ray diffraction patterns of the powder mixtures milled at (a) 500 rpm and (b) 200 rpm for various times, respectively. various times, respectively.

The variation of Dy2O3 grain size with ball-milling time at the rotational speeds of 500 rpm and 200The rpm variation is shown of in Dy Figure2O3 grain 2. Actually, size with the ball-millingsize of Dy2O time3 grain at was the rotationalcalculated speedsfor Dy2O of3 500with rpm cubic and 200structure, rpm is shown not for in FigureDy2O32 .with Actually, monoclinic the size ofstructure, Dy 2O3 grainbecause was ball calculated milling forinduced Dy2O 3 awith phase cubic structure,transformation not for Dy in2 ODy3 with2O3 monoclinicfrom cubic structure,to monoclinic because and ball simultaneously milling induced afrom phase monoclinic transformation to in Dyamorphous.2O3 from It cubic is difficult to monoclinic to calculate and the simultaneously grain size of Dy from2O3 monoclinicwith monoclinic to amorphous. structure. It It can is difficult be seen to calculatethat ball the milling grain results size of in Dy a2 Ofast3 with decrease monoclinic of Dy2O structure.3 grain size It in can the be initial seen stage that ball at both milling 500 resultsrpm and in a fast200 decrease rpm. The of Dyrefinement2O3 grain rate size of in cr theystallite initial size stage is roughly at both 500logarithmic rpm and with 200 rpm.ball-milling The refinement time at 200 rate of crystalliterpm. After size96 h isof roughlyball milling, logarithmic the size of with Dy2O ball-milling3 grain is uptime to approximately at 200 rpm. After60 nm. 96 However, h of ball at milling, 500 therpm, size the of Dydiffraction2O3 grain peaks is up of to Dy approximately2O3 with cubic 60 structure nm. However, are hardly at 500observed rpm, thein the diffraction XRD spectrum peaks of Dyafter2O3 withmilling cubic for structure24 h as shown are hardly in Figure observed 1a, especially, in the XRD when spectrum the ball-milling after milling time for is 24 over h as 48 shown h. Therefore, the size of Dy2O3 with cubic structure is calculated only before 12 h of ball-milling time. It in Figure1a, especially, when the ball-milling time is over 48 h. Therefore, the size of Dy 2O3 with cubic structurecan be seen is calculated that the grain only size before is quickly 12 h of decreased. ball-milling In time.addition, It can after be same seen ball-milling that the grain time, size the is grain quickly size of Dy2O3 phase at the 500 rpm is obviously smaller than that at the 200 rpm. The effect of ball decreased. In addition, after same ball-milling time, the grain size of Dy2O3 phase at the 500 rpm is obviouslymilling on smaller the grain than refinement that at the of 200 powder rpm. Themixtures effect at of 500 ball rpm milling is significantly on the grain more refinement intense than of powder that at 200 rpm. The size of Dy2O3 grain in the powder mixtures after milling for 4 h at 500 rpm is about mixtures at 500 rpm is significantly more intense than that at 200 rpm. The size of Dy2O3 grain in 52 nm, which is smaller than that after milling for 96 h at 200 rpm (approximately 60 nm). the powder mixtures after milling for 4 h at 500 rpm is about 52 nm, which is smaller than that after milling for 96 h at 200 rpm (approximately 60 nm). The morphology evolution of Dy2O3-TiO2 powder mixtures with increasing ball-milling time at 500 rpm is shown in Figure3. Both TiO 2 and Dy2O3 are brittle components, which are fragmented during ball milling and particle size reduces continuously as a consequence of the energy provided during ball milling. The morphology of large particles is changed significantly due to fracture, agglomeration, and deagglomeration processes. The morphology of the original powder mixtures consists of large-sized Dy2O3 particles in micrometer size and small-sized TiO2 particles in nanometer. The shape of the powder particles is irregular. The line-scanning results of elemental Dy, Ti, and O in the characteristic position in Figure3a are shown in Figure3b and also inserted in Figure3a. It can be seen that the small-sized particles are TiO2 component and the large-sized particles are Dy2O3

Materials 2017, 10, 19 5 of 12

components from the variation of the elemental diffraction intensity. The brittle Dy2O3 particles are fragmented by ball-powder-ball collisions, leading to a considerable reduction in the powder particle size and subsequent amorphization as milling time increases. After ball milling for 4 h, the size of particles decreases significantly. A large number of small size of Dy2O3 particles in nanometer can be observed in the milled powders as shown in Figure3c. The morphology of the powder mixtures is transformed to uniform, as shown in Figure3c–g, where the ball-milling time ranges from 4 to 96 h. The morphologies demonstrate that the refining effects of the powder particles are proportional to the ball-milling time for the same rotational speed. After 96 h of ball milling, a large number of nanoparticles agglomerate to form a large-sized particle, as shown in Figure3g. In addition, TiO 2 particles disappear after ball milling for 96 h, as shown in Figure1a. The surfaces of the Dy 2O3 particles in Figure3g are clean compared with those in Figure3a. It can be concluded that the particle size in the powder mixtures is refined to the nanoscale after ball milling for 96 h. The line-scanning results of elemental Dy, Ti, and O in the characteristic position in Figure3g is shown in Figure3h and also inserted in Figure3g, which indicates these elements are uniformly distributed in the ball-milled particles according the variation of the elemental diffraction intensity. In addition, the morphology evolution of powder mixtures at an 200 rpm dose are not provided in the present work because the powder mixtures are only fined and homogenized according to the XRD results as shown in Figure1b. The morphology of mixtures after ball milling for 96 h at 200 rpm is similar with that after ball milling forMaterials 4 h at 5002017, rpm.10, 19 5 of 13

Figure 2. Curves of Dy2O3 grain size vs. ball-milling time. Figure 2. Curves of Dy2O3 grain size vs. ball-milling time. Materials 2017, 10, 19 6 of 13

The morphology evolution of Dy2O3-TiO2 powder mixtures with increasing ball-milling time at 500 rpm is shown in Figure 3. Both TiO2 and Dy2O3 are brittle components, which are fragmented during ball milling and particle size reduces continuously as a consequence of the energy provided during ball milling. The morphology of large particles is changed significantly due to fracture, agglomeration, and deagglomeration processes. The morphology of the original powder mixtures consists of large-sized Dy2O3 particles in micrometer size and small-sized TiO2 particles in nanometer. The shape of the powder particles is irregular. The line-scanning results of elemental Dy, Ti, and O in the characteristic position in Figure 3a are shown in Figure 3b and also inserted in Figure 3a. It can be seen that the small-sized particles are TiO2 component and the large-sized particles are Dy2O3 components from the variation of the elemental diffraction intensity. The brittle Dy2O3 particles are fragmented by ball-powder-ball collisions, leading to a considerable reduction in the powder particle size and subsequent amorphization as milling time increases. After ball milling for 4 h, the size of particles decreases significantly. A large number of small size of Dy2O3 particles in nanometer can be observed in the milled powders as shown in Figure 3c. The morphology of the powder mixtures is transformedFigure 3. to SEM uniform, analysis as results shown of in the Figure powder 3c–g, mixtures where milled the ball-milling at 500 rpm fortime different ranges ball-milling from 4 to 96 h. Figure 3. SEM analysis results of the powder mixtures milled at 500 rpm for different ball-milling The morphologiestimes. The images demonstrate showing the that morphology the refining of ball-milled effects of mixtures the powder for (a )particles 0 h; (c) 4 h;are (d proportional)12 h; (e) 24 to times. The images showing the morphology of ball-milled mixtures for (a) 0 h; (c) 4 h; (d)12 h; (e) 24 h; the ball-millingh; (f) 48 h; and time (g) 96for h; the(b,h same) are EDS rotational analysis speed.results ofAfter ball-milled 96 h ofparticles ball milling, in (a,g), respectively.a large number of (f) 48 h; and (g) 96 h; (b,h) are EDS analysis results of ball-milled particles in (a,g), respectively. nanoparticles agglomerate to form a large-sized particle, as shown in Figure 3g. In addition, TiO2 particlesFigure disappear 4 shows after TEM ball images milling and for corresponding96 h, as shown selected in Figure area 1a. electron The surfaces diffraction of the (SAED) Dy2O3 particlespatterns inof FigureDy2O3-TiO 3g are2 powder clean compared mixtures with milled those for in 4 Figureh and 963a. h.It canAfter be millingconcluded for that4 h, thenano-sized particle sizeball-milled in the powder powder mixtures particles isaggregate refined to to the form nanoscale large-sized after particles, ball milling as shownfor 96 h.in TheFigure line-scanning 4a, due to resultsthe high of activeelemental surface Dy, Ti,energy and Ocreated in the characteristicupon ball milling. position The in size Figure of the3g is original shown inTiO Figure2 particles 3h and is alsoapproximately inserted in 50Figure nm. From3g, which the XRD indicates results, these it can elements be observed are uniformly that ball distributedmilling leads in tothe TiO ball-milled2 particle particlesrefinement, according dissolution, the variation and finally of disappearancethe elemental diffractionafter 96 h. Therefore,intensity. In the addition, main particles the morphology presented evolutionin the TEM of image powder are mixtures Dy2O3. The at anbright 200 zonesrpm dose near arethe notedge provided of the powder in the particlespresent work are the because thin areas the powderwhere the mixtures electron are beam only penetrates. fined and homogenized The dark areas according in the images to the XRDof the results powder as shownparticles in areFigure the 1b.thick The areas morphology where the ofelectron mixtures beam after rarely ball penetrates. milling for Indexing 96 h at 200 and rpm analyzing is similar the withring-shaped that after SAED ball millingpattern fortaken 4 h fromat 500 the rpm. area denoted by the letter “A” indicates that the Dy2O3 grains are already nanocrystalline. The diffraction spots coming from cubic Dy2O3, monoclinic Dy2O3, and TiO2 grains are present in the SAED pattern in Figure 4b. The diffraction halo in the SAED pattern also indicates the formation of an amorphous phase during high-energy ball milling. After ball milling for 96 h, the ball-milled powders agglomerate to form large particles with large thicknesses such that the electron beam rarely penetrates, showing as a dark color in the TEM image in Figure 4c. It is difficult to observe the microstructure of the agglomerated particles. The SAED pattern from the area marked with the letter “B” indicates that the powder mixtures are almost completely converted to the amorphous phase, although sporadic diffraction spots are also present in this pattern. The atom arrangement in the amorphous phase is disordered over long distances, but ordered over short distances. As grain size decreases, the number of atoms at the grain boundaries increases. The proportion of atoms in the crystal volume relative to the crystal boundary decreases. Schwarz and Koch noted that the formation of an amorphous phase in as-milled powders was similar to the amorphization that occurs during the isothermal annealing of crystalline metallic thin films [24]. In high-energy ball milling, the intense deformation accelerates interdiffusion, and the large defect density increases the free energy of the components in the mixture to form an amorphous product.

Materials 2017, 10, 19 6 of 12

Figure4 shows TEM images and corresponding selected area electron diffraction (SAED) patterns of Dy2O3-TiO2 powder mixtures milled for 4 h and 96 h. After milling for 4 h, nano-sized ball-milled powder particles aggregate to form large-sized particles, as shown in Figure4a, due to the high active surface energy created upon ball milling. The size of the original TiO2 particles is approximately 50 nm. From the XRD results, it can be observed that ball milling leads to TiO2 particle reﬁnement, dissolution, and ﬁnally disappearance after 96 h. Therefore, the main particles presented in the TEM image are Dy2O3. The bright zones near the edge of the powder particles are the thin areas where the electron beam penetrates. The dark areas in the images of the powder particles are the thick areas where the electron beam rarely penetrates. Indexing and analyzing the ring-shaped SAED pattern taken from the area denoted by the letter “A” indicates that the Dy2O3 grains are already nanocrystalline. The diffraction spots coming from cubic Dy2O3, monoclinic Dy2O3, and TiO2 grains are present in the SAED pattern in Figure4b. The diffraction halo in the SAED pattern also indicates the formation of an amorphousMaterials 2017, phase10, 19 during high-energy ball milling. 7 of 13

Figure 4. TEM analysis results of the mixtures ball milled at 500 rpm: (a) the bright ﬁeldfield TEM image and ( b) corresponding corresponding SAED SAED pattern pattern of of mixtures mixtures milled milled for for 4 h; 4 h;(c) (cBright) Bright field ﬁeld TEM TEM image image and and (d) (correspondingd) corresponding SAED SAED pattern pattern of mixtures of mixtures milled milled for for96 h. 96 h.

Monoclinic Ln2O3 is initially formed in the mechanical alloying of lanthanum titanate or After ball milling for 96 h, the ball-milled powders agglomerate to form large particles with large dititanate. In Moreno’s research, the formation of Gd2(Ti(1−y)Zry)2O7 pyrochlores occurred in the final thicknesses such that the electron beam rarely penetrates, showing as a dark color in the TEM image step of ball milling starting from an amorphous matrix of Gd2O3, TiO2, and ZrO2 [23]. However, in in Figure4c. It is difficult to observe the microstructure of the agglomerated particles. The SAED the present work, after 96 h of ball milling, the monoclinic Dy2O3 phase could not be detected and pattern from the area marked with the letter “B” indicates that the powder mixtures are almost was completely transformed to the amorphous phase. Moreover, dysprosium titanate also could not completely converted to the amorphous phase, although sporadic diffraction spots are also present be detected. To investigate the sintering behavior of the ball-milled powder mixtures, DSC was in this pattern. The atom arrangement in the amorphous phase is disordered over long distances, carried out on the powder mixtures milled for 96 h at test temperatures ranging from 200 to 1200 °C. but ordered over short distances. As grain size decreases, the number of atoms at the grain boundaries The DSC curve in Figure 5 shows one exothermic peak close to 880 °C and one endothermic peak increases. The proportion of atoms in the crystal volume relative to the crystal boundary decreases. close to 1145 °C. An exothermic peak in a DSC curve can generally be attributed to a transition from Schwarz and Koch noted that the formation of an amorphous phase in as-milled powders was similar disordered to ordered, the recrystallization of original components from the amorphous phase or the to the amorphization that occurs during the isothermal annealing of crystalline metallic thin films [24]. formation of a new compound from the ball-milled amorphous powders. Therefore, subsequent X- In high-energy ball milling, the intense deformation accelerates interdiffusion, and the large defect ray analysis of the ball-milled powders annealed at different temperatures is used to further analyze density increases the free energy of the components in the mixture to form an amorphous product. occurrences in the heating process in more detail. The powder mixtures transform completely to the Monoclinic Ln O is initially formed in the mechanical alloying of lanthanum titanate or dititanate. amorphous state after2 3 ball milling for 96 h. Therefore, it seems feasible that the transition from In Moreno’s research, the formation of Gd (Ti Zr ) O pyrochlores occurred in the final step of ball disordered to ordered produced the exothermic2 (1−y )peaky 2 in7 the DSC curve. In fact, only the amorphous peak is observed in the XRD pattern of the 96 h ball-milled powder after annealing for 24 h at 700 °C. No diffraction peaks for Dy2O3 or TiO2 are detected. Even with prolonged annealing time, the XRD results are the same. Therefore, the exothermic peak in the DSC curve should be related to the new phase generated from the amorphous mixtures.

Materials 2017, 10, 19 7 of 12

milling starting from an amorphous matrix of Gd2O3, TiO2, and ZrO2 [23]. However, in the present work, after 96 h of ball milling, the monoclinic Dy2O3 phase could not be detected and was completely transformed to the amorphous phase. Moreover, dysprosium titanate also could not be detected. To investigate the sintering behavior of the ball-milled powder mixtures, DSC was carried out on the powder mixtures milled for 96 h at test temperatures ranging from 200 to 1200 ◦C. The DSC curve in Figure5 shows one exothermic peak close to 880 ◦C and one endothermic peak close to 1145 ◦C. An exothermic peak in a DSC curve can generally be attributed to a transition from disordered to ordered, the recrystallization of original components from the amorphous phase or the formation of a new compound from the ball-milled amorphous powders. Therefore, subsequent X-ray analysis of the ball-milled powders annealed at different temperatures is used to further analyze occurrences in the heating process in more detail. The powder mixtures transform completely to the amorphous state after ball milling for 96 h. Therefore, it seems feasible that the transition from disordered to ordered produced the exothermic peak in the DSC curve. In fact, only the amorphous peak is observed in the XRD pattern of the 96 h ball-milled powder after annealing for 24 h at 700 ◦C. No diffraction peaks for Dy2O3 or TiO2 are detected. Even with prolonged annealing time, the XRD results are the same. Therefore, the exothermic peak in the DSC curve should be related to the new phase generated from theMaterials amorphous 2017, 10, 19 mixtures. 8 of 13

Figure 5. DSC curve of Dy O -TiO powder mixtures milled for 96 h. Figure 5. DSC curve of Dy22O33-TiO2 powder mixtures milled for 96 h.

TheThe powder mixtures mixtures milled milled for for 96 96h were h were annealed annealed for 3 for h at 3 800, h at 900, 800, 1000, 900, 1050, 1000, 1100, 1050, and 1100, 1150 and °C. 1150The XRD◦C. The results XRD of results the annealed of the annealed powder powder mixtures mixtures are shown are shownin Figure in Figure6a,b. Several6a,b. Several diffraction diffraction peaks peaksdifferent different from the from diffraction the diffraction peaks peaksof Dy2 ofO3 Dy (cubic2O3 (cubicand monoclinic and monoclinic crystal crystal structure) structure) and TiO and2 (rutile TiO2 (rutilestructure) structure) are observed, are observed, which which indicates indicates new compon new componentsents with withcrystal crystal structures structures generated generated from from the theamorphous amorphous mixtures. mixtures. This This experimental experimental phenomenon phenomenon of of synthesizing synthesizing new new co compoundsmpounds is is similar to Stubiˇcar’sStubičar’s researchresearch in in which which the orthorhombicthe orthorhombic ZrTiO 4ZrTiOphase4 wasphase generated was generated from the high-temperaturefrom the high- annealingtemperature of amorphousannealing of mixtures amorphous formed mixtures from formed the ball from milling theof ball aZrO milling2-TiO of2 apowder ZrO2-TiO system2 powder [21] andsystem consistent [21] and with consistent Khor’s with results Khor’s that results zirconia that was zirconia produced was fromproduced annealing from annealing amorphous amorphous mixtures formedmixtures from formed the ballfrom milling the ball of milling an equimolar of an equimolar ZrSiO4 and ZrSiO Al2O4 3andpowder Al2O3system powder [25 system]. [25]. ◦ AccordingAccording toto thethe XRDXRD standardstandard database,database, cubiccubic DyDy22TiOTiO55 {111}{111} presentspresents atat 22θθ == 30.04330.043°,, hexagonal ◦ ◦ DyDy22TiO5 {102}{102} presents presents at at 2θ 2θ = =32.411°, 32.411 orthorhombic, orthorhombic Dy Dy2TiO2TiO5 {201}5 {201} presents presents at 2 atθ = 2 θ29.554°,= 29.554 and, ◦ andpyrochlore pyrochlore Dy2Ti Dy2O2Ti7 {111}2O7 {111} presents presents at 2θ at = 2 θ30.698°;= 30.698 the; thedifference difference in the in the above above diffraction diffraction angles angles is isnot not very very large. large. Three Three diffraction diffraction peaks peaks are are observed observed in in the the XRD XRD patterns patterns of the powderpowder mixturesmixtures annealedannealed forfor 33 hh at 800, 900, and 1000 °C.◦C. The main diffraction peak representing the crystalline phase is at 2θ = 30.0°, as shown in Figure 6b. Therefore, the newly formed product in the powder system annealed between 800 °C and 1000 °C is not pyrochlore Dy2Ti2O7. This result is different than that found in Amit Sinha’s research in which pyrochlore Dy2Ti2O7 was initially created in the formation of Dy2TiO5 [19]. In their research using an equimolar Dy2O3-TiO2 system, the chemical reaction of Dy2O3 and TiO2 initially formed pyrochlore Dy2Ti2O7, and then Dy2Ti2O7 reacted with the remaining Dy2O3 to form orthorhombic Dy2TiO5. According to the equilibrium phase diagram, the Dy2TiO5 phase has three crystal structure types 1350 ℃ 1680 ℃ depending on the temperature: orthorhombic ርۛۛሮ hexagonal ርۛۛሮ cubic [15]. Transformation from the high-temperature phase to the low-temperature phase is an exothermic process. For example, the polymorphic transformation of Gd2TiO5 from hexagonal to orthorhombic produced an exothermic peak in the DSC curve [7]. However, in the present work, there is only one exothermic peak in the DSC curve. Although the diffraction peaks in the XRD patterns match well with the diffraction peaks in the standard pattern of cubic Dy2TiO5, the generated phase should not be cubic Dy2TiO5 because low-temperature phases of orthorhombic and hexagonal Dy2TiO5 were not detected after annealing over temperatures ranging from 700 to 1000 °C for annealing times ranging from several minutes to 3 h; additionally, as mentioned above, the formation temperature of cubic Dy2TiO5 is over 1680 °C. It is not possible to achieve such a high temperature during ball milling. Therefore, the high-temperature phase of cubic Dy2TiO5 should not be produced. Instead, the generated phase is an intermediate phase that has a similar crystal structure to cubic Dy2TiO5 and is a metastable state that phase transforms to orthorhombic Dy2TiO5 when the annealing temperature is above 1050 °C.

Materials 2017, 10, 19 9 of 13

After annealing for 3 h at 1100 °C, orthorhombic Dy2TiO5 and a small amount of pyrochlore Dy2Ti2O7 and cubic Dy2O3 are detected; notably, cubic Dy2TiO5 is not observed in the ball-milled powder mixtures. The change of the grain size of the main characteristic phase in the annealed powder mixtures

Materialswith annealing2017, 10, 19 temperature is shown in Figure 6c. According to the above results, the grain size8 of of 12 the intermediate phase and orthorhombic Dy2TiO5 are calculated using Suryanarayana and Grant Norton’s formula. It can be seen that the grain size increases with increasing annealing temperature. ◦ isThe at grain 2θ = 30.0size ,of as the shown intermediate in Figure phase6b. Therefore, is about 27 the nm newly in the formed powder product mixtures in the annealed powder for system 3 h at ◦ ◦ annealed800 °C and between is about 800 275C andnm 1000in theC powder is not pyrochlore mixtures annealed Dy2Ti2O7 .for This 3 resulth at 1050 is different °C. Because than thatthe foundintermediate in Amit phase Sinha’s transforms research into whichorthorhombic pyrochlore Dy2 DyTiO25Ti when2O7 was the initiallyannealing created temperature in the formation is above of1050 Dy °C,2TiO the5 [ 19grain]. In size their of researchorthorhombic using Dy an2 equimolarTiO5 is calculated Dy2O3-TiO at the2 system, annealing the temperature chemical reaction ranging of Dyfrom2O 31050and °C TiO to2 initially 1150 °C. formed The grain pyrochlore size of Dy orthorhombic2Ti2O7, and then Dy Dy2TiO2Ti5 2isO 7aboutreacted 43 withand the380 remaining nm after Dyannealing2O3 to form for 3 orthorhombich at 1050 °C and Dy 21150TiO5 °C,. respectively.

Figure 6. XRD patternspatterns ofof thetheball-milled ball-milled powder powder mixtures mixtures annealed annealed for for 3 h3 ath at various various temperatures temperatures at diffractionat diffraction angles angles 2θ ranging2θ ranging from from (a) 20 (a◦) to20° 65 to◦ and65° and (b) 28 (b◦)to 28° 31 to◦;( 31°;c) Curves (c) Curves of the of grain the grain size of size main of characteristicmain characteristic phase phase vs. annealing vs. annealing temperature. temperatur The powdere. The powder mixtures mixtures were previously were previously milled for milled 96 h atfor 500 96 rpm.h at 500 rpm.

In addition, the pressure created during ball milling is not high enough to transform the crystal According to the equilibrium phase diagram, the Dy2TiO5 phase has three crystal structure types 2 5 structure of Dy TiO from orthorhombic to 1350hexagonal◦C or from1680 ◦hexagonalC to cubic. The average dependingpressure on on the the contact temperature: surface orthorhombicof two colliding↔ millhexagonal balls is approximately↔ cubic [8.515]. GPa Transformation [26], which is from far thebelow high-temperature the 100 GPa needed phase to induce to the low-temperature a pressure wave to phase cause is the an Gd exothermic2TiO5 phase process. transformation For example, from thethe polymorphiclow-temperature transformation orthorhombic of Gd phase2TiO 5tofrom the hi hexagonalgh-temperature to orthorhombic hexagonal produced phase [27]. an exothermicTherefore, peakthe intermediate in the DSC phase curve is [7 ].not However, produced in by the collision present pressure. work, there Further is only investigation one exothermic is needed peak into in thethe DSCcause curve. of the Althoughformation the of the diffraction intermediate peaks phase. in the XRD patterns match well with the diffraction peaks in theTo standard further patterninvestigate of cubic the annealing Dy2TiO5, thebehavior generated of the phase ball-milled should powder not be cubic mixtures, Dy2TiO the5 becausepowder low-temperaturemixtures milled for phases 96 h ofat orthorhombic200 rpm were andsintered hexagonal for 3 h Dy at2 800,TiO5 900,were 1000, not detected1050, 1100, after and annealing 1150 °C. ◦ overThe phase temperatures evolution ranging of the annealed from 700 powder to 1000 mixturesC for annealing identified times by XRD ranging analysis from is severalshown in minutes Figure to7. ◦ 3Under h; additionally, these ball-milling as mentioned conditions, above, thethe formation Dy2O3-TiO temperature2 powder mixtures of cubic Dy are2TiO homogenized,5 is over 1680 andC. Itthe is notpolymorphic possible to transf achieveormation such a of high Dy temperature2O3 from cubic during to monoclinic ball milling. is not Therefore, observed the in high-temperature Figure 7a. In the phaseXRD pattern of cubic of Dy the2TiO powder5 should mixtures not be produced. annealed Instead,for 3 h at the 1000 generated °C, diffraction phase is anpeaks intermediate for Dy2O phase3 and thatDy2Ti has2O7 a phase similar are crystal observed. structure However, to cubic the Dy main2TiO diffraction5 and is a metastable peak for orthorhombic state that phase Dy2 transformsTiO5 is not ◦ todetected. orthorhombic Under Dythese2TiO conditions,5 when the the annealing annealed temperature powder mixtures is above are 1050 composedC. After of annealingcubic Dy2O for3 and 3 h ◦ atpyrochlore 1100 C, orthorhombicDy2Ti2O7. The Dypowder2TiO5 mixturesand a small generate amount the of orthorhombic pyrochlore Dy Dy2Ti2TiO2O75 andphase cubic at 1050 Dy2 °CO3 andare detected;are composed notably, of cubic cubic Dy Dy2O2TiO3, pyrochlore5 is not observed Dy2Ti2O in7, and the ball-milledorthorhombic powder Dy2TiO mixtures.5. The powder mixtures are almostThe change completely of the transformed grain size of to the orthorhombic main characteristic Dy2TiO phase5 afterin annealing the annealed at 1150 powder °C for mixtures 3 h. The withdiffraction annealing peak temperature intensity of is orthorhombic shown in Figure Dy6c.2TiO According5 gradually to the increases above results, with increasing the grain sizeannealing of the intermediatetemperature, phase and simultaneously, and orthorhombic the Dydiffraction2TiO5 are peak calculated intensity using of Dy Suryanarayana2O3 and Dy2Ti and2O7 Grantdecreases Norton’s with formula.increasing It canannealing be seen temperature. that the grain For size increasespowder mi withxtures increasing ball milled annealing for temperature.96 h at 200 Therpm, grain the ◦ sizeevolutionary of the intermediate behavior at phase various is aboutannealing 27 nm temperatures in the powder is consistent mixtures annealedwith the data for 3presented h at 800 inC andRef. ◦ is[19], about in which 275 nm Amit in the Sinha powder reported mixtures that annealed the chemical for 3 hreaction at 1050 ofC. Dy Because2O3 and the TiO intermediate2 initially formed phase ◦ transformspyrochlore toDy orthorhombic2Ti2O7, and then Dy2 TiODy25Tiwhen2O7 reacted the annealing with the temperature remaining is Dy above2O3 to 1050 formC, orthorhombic the grain size ◦ ◦ of orthorhombic Dy2TiO5 is calculated at the annealing temperature ranging from 1050 C to 1150 C. ◦ The grain size of orthorhombic Dy2TiO5 is about 43 and 380 nm after annealing for 3 h at 1050 C and 1150 ◦C, respectively. In addition, the pressure created during ball milling is not high enough to transform the crystal structure of Dy2TiO5 from orthorhombic to hexagonal or from hexagonal to cubic. The average Materials 2017, 10, 19 9 of 12 pressure on the contact surface of two colliding mill balls is approximately 8.5 GPa [26], which is far below the 100 GPa needed to induce a pressure wave to cause the Gd2TiO5 phase transformation from the low-temperature orthorhombic phase to the high-temperature hexagonal phase [27]. Therefore, the intermediate phase is not produced by collision pressure. Further investigation is needed into the cause of the formation of the intermediate phase. To further investigate the annealing behavior of the ball-milled powder mixtures, the powder mixtures milled for 96 h at 200 rpm were sintered for 3 h at 800, 900, 1000, 1050, 1100, and 1150 ◦C. The phase evolution of the annealed powder mixtures identiﬁed by XRD analysis is shown in Figure7. Under these ball-milling conditions, the Dy2O3-TiO2 powder mixtures are homogenized, and the polymorphic transformation of Dy2O3 from cubic to monoclinic is not observed in Figure7a. In the ◦ XRD pattern of the powder mixtures annealed for 3 h at 1000 C, diffraction peaks for Dy2O3 and Dy2Ti2O7 phase are observed. However, the main diffraction peak for orthorhombic Dy2TiO5 is not detected.Materials 2017 Under, 10, 19 these conditions, the annealed powder mixtures are composed of cubic Dy2O103 andof 13 ◦ pyrochlore Dy2Ti2O7. The powder mixtures generate the orthorhombic Dy2TiO5 phase at 1050 C andDy2TiO are5 composed. However, of this cubic experimental Dy2O3, pyrochlore phenomenon Dy2Ti is2 Odifferent7, and orthorhombic from that observed Dy2TiO in5. the The annealed powder ◦ mixturespowder mixtures are almost milled completely for 96 h at transformed 500 rpm, as to shown orthorhombic in Figure Dy6, due2TiO to5 afterthe initial annealing conditions at 1150 of theC forball-milling 3 h. The diffractionmixtures. peak intensity of orthorhombic Dy2TiO5 gradually increases with increasing annealingThe change temperature, of the grain and simultaneously,size of main characteristic the diffraction phase peak in the intensity annealed of powder Dy2O3 andmixtures Dy2Ti with2O7 decreasesannealing withtemperature increasing is annealingshown in Figure temperature. 7c. The For grain powder size of mixtures the cubic ball Dy milled2O3, pyrochlore for 96 h at Dy 2002Ti rpm,2O7 theand evolutionary orthorhombic behavior Dy2TiO at5 are various calculated annealing using temperatures Suryanarayana is consistent and Grant with Norton’s the data formula. presented The in Ref.grain [19 size], in of which cubic AmitDy2O Sinha3 and pyrochlore reported that Dy the2Ti2 chemicalO7 increases reaction with ofincreasing Dy2O3 and annealing TiO2 initially temperature. formed pyrochloreBecause the Dy orthorhombic2Ti2O7, and thenDy2TiO Dy52 isTi 2detectedO7 reacted in the with powder the remaining mixtures Dy annealed2O3 to form at 1050 orthorhombic °C for 3 h, Dythe2 grainTiO5. size However, of orthorhombic this experimental Dy2TiO5 phenomenonis calculated when is different the annealin from thatg temperature observed inis theover annealed 1050 °C. powderThe grain mixtures size of orthorhombic milled for 96 hDy at2TiO 5005 rpm, is about as shown 38 nm inand Figure 395 nm6, due after to annealing the initial 3 conditions h at 1050 °C of and the ball-milling1150 °C, respectively. mixtures.

Figure 7. XRD patternspatterns ofof thethe ball-milledball-milled powder powder mixtures mixtures annealed annealed for for 3 h3 h at at various various temperatures temperatures at diffractionat diffraction angles angles 2θ 2rangingθ ranging from from (a )(a 20) 20°◦ to to 65 65°◦ and and (b ()b 28) 28°◦ to to 31.0 31.0°;◦;( (cc)) CurvesCurves ofof thethe graingrain sizesize of characteristic phase vs. annealing temperature. The powder mixtures were previously milled for 96 h at 200 rpm.

Figure 8a,b shows the bright field TEM image and corresponding SAED pattern of the ball- The change of the grain size of main characteristic phase in the annealed powder mixtures with milled Dy2O3-TiO2 powder mixtures annealed at 1000 °C for 3 h, respectively. The powder mixtures annealing temperature is shown in Figure7c. The grain size of the cubic Dy O , pyrochlore Dy Ti O are previously milled for 96 h at 500 rpm. After annealing for 3 h, the grain2 size3 of powder mixtures2 2 7 and orthorhombic Dy TiO are calculated using Suryanarayana and Grant Norton’s formula. The grain is kept in nanometer 2scale,5 which can be supported by the corresponding SAED pattern taken from size of cubic Dy2O3 and pyrochlore Dy2Ti2O7 increases with increasing annealing temperature. Because the region marked letter “A” in Figure 8a. The diffraction ring is a typical SAED◦ pattern of the orthorhombic Dy2TiO5 is detected in the powder mixtures annealed at 1050 C for 3 h, the nanocrystal materials. After analyzing and indexing the ring-shaped SAED pattern, it is indicated◦ grain size of orthorhombic Dy2TiO5 is calculated when the annealing temperature is over 1050 C. that this SAED pattern belongs to the intermediate Dy2TiO5 phase that has a similar crystal structure◦ The grain size of orthorhombic Dy2TiO5 is about 38 nm and 395 nm after annealing 3 h at 1050 C and to cubic Dy2TiO5, which is in accord with the XRD results as shown in Figure 6. The small-sized 1150 ◦C, respectively. powders agglomerate to form large particles with large thicknesses as shown in Figure 8a. The bright Figure8a,b shows the bright ﬁeld TEM image and corresponding SAED pattern of the ball-milled zones near the edge of the powder particles is the thin area where the electron beam penetrates. The Dy O -TiO powder mixtures annealed at 1000 ◦C for 3 h, respectively. The powder mixtures are dark2 area3 in2 the image of the powder particles is the thick area where the electron beam rarely penetrates. After annealing, the amorphous ball-milled powder mixtures are changed to the intermediate Dy2TiO5 phase with crystal structure. Figure 8c,d is the bright field TEM image and corresponding SAED pattern of the annealed Dy2O3-TiO2 powder mixtures that were previously milled for 96 h at 200 rpm. Indexing and analyzing the ring-shaped SAED pattern taken from the area denoted by the letter “B” indicates that the particles are composed of cubic Dy2O3 and pyrochlore Dy2Ti2O7, which is accord with the XRD results of the powder mixtures annealed at 1000 °C for 3 h as shown in Figure 7. This experimental result is different from that observed in the annealed powder mixtures milled for 96 h at 500 rpm.

Materials 2017, 10, 19 10 of 12 previously milled for 96 h at 500 rpm. After annealing for 3 h, the grain size of powder mixtures is kept in nanometer scale, which can be supported by the corresponding SAED pattern taken from the region marked letter “A” in Figure8a. The diffraction ring is a typical SAED pattern of nanocrystal materials. After analyzing and indexing the ring-shaped SAED pattern, it is indicated that this SAED pattern belongs to the intermediate Dy2TiO5 phase that has a similar crystal structure to cubic Dy2TiO5, which is in accord with the XRD results as shown in Figure6. The small-sized powders agglomerate to form large particles with large thicknesses as shown in Figure8a. The bright zones near the edge of the powder particles is the thin area where the electron beam penetrates. The dark area in the image of the powder particles is the thick area where the electron beam rarely penetrates. After annealing, the amorphous ball-milled powder mixtures are changed to the intermediate Dy2TiO5 phase with crystal structure. Figure8c,d is the bright ﬁeld TEM image and corresponding SAED pattern of the annealed Dy2O3-TiO2 powder mixtures that were previously milled for 96 h at 200 rpm. Indexing and analyzing the ring-shaped SAED pattern taken from the area denoted by the letter “B” indicates that the particles are composed of cubic Dy2O3 and pyrochlore Dy2Ti2O7, which is accord with the XRD results of the powder mixtures annealed at 1000 ◦C for 3 h as shown in Figure7. This experimental result is different fromMaterials that 2017 observed, 10, 19 in the annealed powder mixtures milled for 96 h at 500 rpm. 11 of 13

Figure 8. The bright field TEM images and corresponding SAED patterns of the ball-milled Dy2O3- Figure 8. The bright ﬁeld TEM images and corresponding SAED patterns of the ball-milled Dy2O3-TiO2 TiO2 powder mixtures annealed at 1000 °C for 3 h, (a,b) the powder mixtures are previously milled powder mixtures annealed at 1000 ◦C for 3 h, (a,b) the powder mixtures are previously milled for 96 h atfor 500 96 h rpm; at 500 (c, drpm;) the ( powderc,d) the powder mixtures mixtures are previously are previously milled for milled 96 h atfor 200 96 rpm.h at 200 rpm.

4. Conclusions 4. Conclusions The microstructural evolution of Dy2O3-TiO2 powder mixtures during ball milling and post- The microstructural evolution of Dy O -TiO powder mixtures during ball milling and milled annealing was investigated using TEM,2 SEM,3 2XRD, and DSC. The conclusions can be made as post-milled annealing was investigated using TEM, SEM, XRD, and DSC. The conclusions can be made follows: as follows: 1. The ball-milling parameters had a great effect on ball milling and the subsequent annealing process. 2. At 500 rpm rotation speeds, the mixtures were fined, homogenized, nanocrystallized, and then completely amorphized, and the crystal structure of Dy2O3 was transformed from cubic to monoclinic. The amorphous transformation resulted from monoclinic Dy2O3, not from cubic Dy2O3. However, at 200 rpm rotation speeds, the Dy2O3-TiO2 powder mixtures were only homogenized, and the polymorphic transformation of Dy2O3 from cubic to monoclinic was not observed. Meanwhile, the powder mixtures did not transform to the amorphous phase. 3. The powder mixtures milled for 96 h at 500 rpm were annealed for 3 h at a temperature range of 800 to 1000 °C. An intermediate phase with a crystal structure similar to that of cubic Dy2TiO5 was synthesized, which was a metastable phase that transformed to orthorhombic Dy2TiO5 when the annealing temperature was above 1050 °C. However, the powder mixtures milled for 96 h at 200 rpm did not transform to the amorphous phase. The annealing behavior showed that the chemical reaction of Dy2O3 with TiO2 initially formed pyrochlore Dy2Ti2O7, and then Dy2Ti2O7 reacted with the remaining Dy2O3 to form orthorhombic Dy2TiO5.

Acknowledgments: The work was supported by the National Natural Science Foundation of China through Grant No. 11305136.

Author Contributions: Guang Ran conceived and designed the experiments; Jinhua Huang performed the experiments and analyzed the data; Qiang Shen conducted the TEM experiment; Jinhua Huang, Jianxin Lin, Penghui Lei, and Xina Wang wrote the paper under the supervision of Guang Ran. All authors contributed to the scientific discussion of the results and reviewed the manuscript.

Materials 2017, 10, 19 11 of 12

1. The ball-milling parameters had a great effect on ball milling and the subsequent annealing process. 2. At 500 rpm rotation speeds, the mixtures were ﬁned, homogenized, nanocrystallized, and then completely amorphized, and the crystal structure of Dy2O3 was transformed from cubic to monoclinic. The amorphous transformation resulted from monoclinic Dy2O3, not from cubic Dy2O3. However, at 200 rpm rotation speeds, the Dy2O3-TiO2 powder mixtures were only homogenized, and the polymorphic transformation of Dy2O3 from cubic to monoclinic was not observed. Meanwhile, the powder mixtures did not transform to the amorphous phase. 3. The powder mixtures milled for 96 h at 500 rpm were annealed for 3 h at a temperature range ◦ of 800 to 1000 C. An intermediate phase with a crystal structure similar to that of cubic Dy2TiO5 was synthesized, which was a metastable phase that transformed to orthorhombic Dy2TiO5 when the annealing temperature was above 1050 ◦C. However, the powder mixtures milled for 96 h at 200 rpm did not transform to the amorphous phase. The annealing behavior showed that the chemical reaction of Dy2O3 with TiO2 initially formed pyrochlore Dy2Ti2O7, and then Dy2Ti2O7 reacted with the remaining Dy2O3 to form orthorhombic Dy2TiO5.

Acknowledgments: The work was supported by the National Natural Science Foundation of China through Grant No. 11305136. Author Contributions: Guang Ran conceived and designed the experiments; Jinhua Huang performed the experiments and analyzed the data; Qiang Shen conducted the TEM experiment; Jinhua Huang, Jianxin Lin, Penghui Lei, and Xina Wang wrote the paper under the supervision of Guang Ran. All authors contributed to the scientific discussion of the results and reviewed the manuscript. Conflicts of Interest: The authors declare no conflicts of interest.

References

1. Ran, G.; Zhou, J.E.; Xi, S.Q.; Li, P.L. Formation of nanocrystalline and amorphous phase of Al-Pb-Si-Sn-Cu powder during mechanical alloying. Mater. Sci. Eng. A 2006, 416, 45–50. 2. Eckert, J.; Schultz, L.; Urban, K. Formation of quasicrystals by mechanical alloying. Appl. Phys. Lett. 1989, 55, 117–119. [CrossRef] 3. Luo, Y.; Ran, G.; Chen, N.J.; Shen, Q.; Zhang, Y.L. Microstructural evolution, thermodynamics and kinetics

of Mo-Tm2O3 powder mixtures during ball milling. Materials 2016, 9, 834. [CrossRef] 4. Kimura, H.; Kimura, M.; Takada, F. Development of an extremely high energy ball mill for solid state amorphizing transformations. J. Less Common Met. 1988, 40, 113–118. [CrossRef] 5. Gilman, P.S.; Benjamin, J.S. Mechanical alloying. Annu. Rev. Mater. Sci. 1983, 13, 279–300. [CrossRef] 6. Kong, L.B.; Ma, J.; Zhu, W.; Tan, O.K. Preparation of PMN-PT ceramics via a high-energy ball milling process. J. Alloys Compd. 2002, 335, 290–296. [CrossRef] 7. G-Martinez, G.; MGonzalez, L.G.; Escalante-García, J.I.; Fuentes, A.F. Phase evolution induced by mechanical

milling in Ln2O3:TiO2 mixtures (Ln = Gd and Dy). Powder Technol. 2005, 152, 72–78. [CrossRef] 8. Fuentes, A.F.; Takacs, L. Preparation of multicomponent oxides by mechanochemical methods. J. Mater. Sci. 2013, 48, 598–611. [CrossRef] 9. Suryanarayana, C. Mechanical alloying and milling. Prog. Mater. Sci. 2001, 46, 1–184. [CrossRef] 10. Khakpour, Z.; Youzbashi, A.A.; Maghsoudipour, A.; Ahmadi, K. Synthesis of nanosized gadolinium doped ceria solid solution by high energy ball milling. Powder Technol. 2011, 214, 117–121. [CrossRef]

11. Huang, J.H.; Ran, G.; Liu, T.J.; Shen, Q.; Li, N. Microstructure and Physical Properties of Tb2TiO5 Neutron Absorber Synthesized by Ball Milling and Sintering. J. Mater. Eng. Perform. 2016, 25, 4266–4273. [CrossRef] 12. Risovany, V.D.; Varlashova, E.E.; Suslov, D.N. Dysprosium titanate as an absorber material for control rods. J. Nucl. Mat. 2000, 281, 84–89. [CrossRef] 13. Risovany, V.D.; Klochkov, E.P.; Varlashova, E.E. Hafnium and dysprosium titanate based control rods for thermal water-cooled reactors. At. Energy 1996, 81, 764–769. [CrossRef] 14. Kermit, W.; Theilacker, J.S. Neutron Absorber Materials for Reactor Control; Naval Reactors Division of Reactor Development United States Atomic Energy Commission: Washington, DC, USA, 1962.

15. Petrova, M.A.; Novikova, A.S.; Grebenshchikov, R.G. Polymorphism of rare earth titanates of Ln2TiO5 composition. Izv. Akad. Nauk SSSR. Neorg. Mater. 1982, 18, 287–291. Materials 2017, 10, 19 12 of 12

16. Jung, C.H.; Kim, C.J.; Lee, S.J. Synthesis and sintering studies on Dy2TiO5 prepared by polymercarrier chemical process. J. Nucl. Mater. 2006, 354, 137–142. [CrossRef] 17. Panneerselvam, G.; Venkata Krishnan, R.; Antony, M.P.; Nagarajam, K.; Vasudevan, T.; Vasuadeva Rao, P.R. Thermophysical measurements on dysprosium and gadolinium titanates. J. Nucl. Mater. 2004, 327, 220–225. [CrossRef] 18. Kim, J.S.; Kim, H.S.; Jeong, C.Y. Phase study on the DyxTiyOz pellets. In Proceedings of the Transactions of the Korean Nuclear Society Autumn Meeting, Busan, Korea, 27–28 October 2005. 19. Sinha, A.; Prakash, B. Development of Dysprosium Titanate Based Ceramics. J. Am. Ceram. Soc. 2005, 88, 1064–1066. [CrossRef] 20. Suryanarayana, C.; Norton, M.G. X-ray Diffraction: A Practical Approach; Plenum Press: New York, NY, USA, 1998. 21. Stubiˇcar, M.; Bermanec, V.; Stubiˇcar, N.; Dudrnovski, D.; Drumes, D. Microstructure evolution of an

equimolar powder mixture of ZrO2-TiO2 during high-energy ball-milling and post-annealing. J. Alloys Compd. 2001, 316, 316–320. [CrossRef]

22. Gajovi´c,A.; Djerdj, I.; Furi´c,K.; Schlögl, R.; Su, D.S. Preparation of nanostructured ZrTiO4 by solid state reaction in equimolar mixture of TiO2 and ZrO2. J. Am. Ceram. Soc. 2006, 89, 2196–2205. [CrossRef] 3+ 3+ 23. Moreno, K.J.; Rodrigo, R.S.; Fuentes, A.F. Direct synthesis of A2(Ti(1−y)Zry)2O7 (A = Gd ,Y ) solid solutions by ball milling constituent oxides. J. Alloys Compd. 2005, 390, 230–235. [CrossRef] 24. Koch, C.C.; Cavin, O.B. Preparation of “amorphous”Ni60Nb40 by mechanical alloying. Appl. Phys. Lett. 1986, 49, 146–148. [CrossRef]

25. Khor, K.A.; Li, Y. Effect of mechanical alloying on the reaction sintering of ZrSiO4 and Al2O3. Mater. Sci. Eng. A 1998, 256, 271–279. [CrossRef] 26. Huang, J.H.; Lu, J.Q.; Ran, G.; Chen, N.J.; Qu, P.D. Formation of nanocrystalline and amorphization phase of

Fe-Dy2O3 powder mixtures induced by ball milling. J. Mater. Res. 2016.[CrossRef] 27. Shcherbakova, L.G.; Kolesnikov, A.V.; Breusov, O.N. Investigation of TiO2-Ln2O3 systems under shock wave action. Inorg. Mater. 1979, 15, 1724–1729.

© 2016 by the authors; licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC-BY) license (http://creativecommons.org/licenses/by/4.0/).