Original Research Papers
Formation of Nanoscale Layered Structures and Subsequent Transformations during Mechanical Alloying of Ni60Nb40 Powder Mixture in a Low Energy Ball Mill
2015 Volume 32 Pages 196-206
Details
2015 Volume 32 Pages 196-206
Elemental Ni-Nb powder mixture containing 40 atom % Nb were mechanically alloyed in a low energy ball mill, allowing slower processing and easier progressive observation at intermediate milling times. The evolution of morphology and structure of the powders were investigated as a function of milling time by transmission and scanning electron microscopy and X-ray diffractometry. The results revealed that an ultrafine Ni/Nb layered structure with a typical thickness of ~30 nm, containing nanoscale size grains with a typical size of ~15 nm, develops in powder particles during mechanical alloying. This microstructure provides numerous high speed diffusion paths such as sub-grain boundaries and dislocation networks, allowing a high diffusion rate at low temperature and therefore permits different solid-state reactions to take place kinetically. Under mechanical alloying conditions used here continued milling led to a fully amorphous structure. Because of non-uniform plastic deformation the kinetic requirements for the amorphization reaction at the edges of particles is satisfied prior to the centres, resulting in an inhomogeneous progress of amorphization reaction from the edges towards the centres of particles.
Mechanical alloying is a solid-state synthesis route for materials using ball milling process. The starting powder, often a mixture of elemental constituent powders, is charged into a suitable mill and processed for several hours depending on experimental conditions as well as desired microstructure and properties. Often, mechanical alloying has been performed in ball mills including planetary ball mills, centrifugal ball mills and vibratory ball mills, but several other types of mills such as rod mills can also be used. The foremost restriction for mechanical alloying is a sufficiently high energy which is provided by the kinetic energy of the balls on impact with the powder particles. The energy of milling machines depends not only on the power supplied to drive the milling chamber, but also on the internal mechanics of the specific mill. The final powder produced by mechanical alloying can be subsequently either consolidated by standard powder metallurgy techniques into bulk materials with desirable properties or deposited on surfaces of engineering parts using various thermal spraying methods.
From the viewpoint of mass transfer the ball milling process can be divided into two major categories: a) Milling processes which involve mass transfer between the components. Such cases can occur during milling of multi-component powders (e.g. a mixture of elemental or dissimilar alloy powders) and are associated with compositional changes of powder particles. This category is known as “mechanical alloying” (MA). Material transfer by diffusion of components during MA is significantly accelerated by lattice defects and by a momentary increase in temperature of particles trapped between colliding balls. Modeling as well as inferences drawn from the product structure suggests modest temperature rises (∼100–200 °C) result from the kinetic energy of the milling balls. However, if a large exothermic reaction is involved much higher local temperatures can be produced by milling (Takacs, 2002). b) Milling processes that involve no mass transfer and change the structure of a single composition powder (e.g. an elemental or intermetallic compound powder). These cases have often been termed “mechanical milling” (MM) or “mechanical grinding” (MG).
In MA experiments not only the milling time but also further parameters such as the type of mill, the milling tool materials, the energy of milling, the ball-to-powder weight ratio, the milling temperature and the milling atmosphere may have significant influence on phase transformations occurring during MA and therefore, on the final structure. Hence, by controlling these parameters a large variety of materials ranging from stable to meta-stable structures can be synthesised by MA.
MA/MM leads to the significant refinement of microstructure, which is often accompanied with transformation to metastable structures. These techniques increase the range of obtainable microstructures and can therefore lead to significant improvements in properties of materials. There are lots of papers published on synthesis and processing of various materials by MA/MA during the past decade showing improved microstructures and properties. In this paper the progressive development of nanoscaled layered structure during MA process is reported in detail in case of Ni-Nb alloy, as a typical ductile-ductile system. The MA of Ni-Nb system using high energy ball mills has been previously reported (Koch et al., 1983; Petzoldt, 1988; Zhang, 2004). However, the MA behaviour has not been followed in detail because of the rapid processing times used in high energy ball mills. In the present study a low energy ball mill was used to allow slower processing and easier progressive observation at intermediate milling times. Principally it is possible to achieve similar results in a high energy ball mill by controlling the milling speed but one should note that the type of microstructure, structural evolution, kinetics and thermodynamics of transformations, etc., depend not only on level of energy of mill machine but also on the internal mechanics of the specific mill (the movement mode of balls in the bowl). In planetary ball mills the most common phenomenon is the impact of balls on powder particles while in a centrifugal ball mill the dominate event is rolling of balls inside the bowl.
99.9 % pure Ni powder and 99.8 % pure Nb powder, from Johnson Matthey, were used as starting materials. Fig. 1 (a) and (b) shows scanning electron microscopy (SEM) images of the elemental Ni and Nb particles.
SEM micrographs of as-received Ni and Nb powder particles and Ni60Nb40 powder particles after different milling times.
As shown, the Ni particles had a nearly uniform size of ∼25 to 50 μm with a morphology consisting of clusters of several small spherical particles. The Nb particles were angular in shape with a wide size distribution from ∼5 to 110 μm. All elemental powders were stored in a vacuum chamber. The powders were weighed and mixed manually to give nominal atomic composition of Ni60Nb40.
Mechanical alloying was carried out using a laboratory Fritsch Pulverisette 6 centrifugal ball mill. Pulverisette 6 is a low energy type mill, capable of achieving up to 8 times gravitational acceleration. Higher energy ball mill systems, as exemplified by a planetary ball mill, are capable of producing a higher acceleration (e.g. 45 times gravitational acceleration). A bowl containing five balls was fitted on a supporting disc. The drive mechanism rotated the supporting disc on a circular trajectory. The milling media consisted of five 20 mm diameter balls confined in an 80 ml volume bowl. The bowl and ball materials were hardened chromium steel. For each MA run 17 g of powder was used. To prevent oxidation or nitridation of the powder, MA was performed under an Ar atmosphere. The powder mixture and bowl were placed into a glove bag. To reduce the oxygen content, the glove bag was flushed out with a high-purity Ar gas and then fully filled. The powder mixture was loaded into the bowl then sealed using a teflon O-ring, covered with a thin layer of sealing grease, and sealing tape. The bowl was removed from the glove bag and mounted on the mill machine. In all MA runs the ball-to-powder weight ratio was about 10:1 and the angular velocity of the supporting disc was approx. 600 rpm. MA was performed nominally at room temperature although a small increase in temperature of the bowl exterior (ΔT≈20 °C), measured using a digital, portable thermometer, was observed during process.
The MA was interrupted after selected time intervals and the bowl was opened in air. A small amount of powder was removed for further characterization and then the bowl was resealed in the Ar-filled glove bag and MA continued.
The structure of MA samples were investigated by X-ray diffraction (XRD) analysis in a Philips vertical X-ray diffractometer (PW1710) with a scanning speed of 0.03° s−1 in a step-counting mode. Ni-filtered CuKα radiation (λ = 0.1542 nm) produced by a generator voltage of 35 kV and a generator current of 50 mA was used.
The grain size of the MA powder was estimated by analysing the X-ray diffraction peak broadening. Since strain in the crystal lattice also contributes to broadening of the XRD peaks, the approach of Williamson and Hall (1953) was used in order to separate the two effects of grain size and strain after correction for the instrumental broadening.
The particle morphologies were investigated in a Hitachi S-530 SEM operating at 25 kV. Powder samples were fixed in small quantities on a sample holder by Ag conductive paint. At least 40 separate particles were chosen for the measurement of powder particle size.
A Philips CM20 transmission electron microscope (TEM) operating at 200 kV was used for observation and characterization of the internal structure of as-milled powders. The sample preparation for TEM consisted of mixing powder with a small amount of a TEM suitable epoxy glue (Gatan G-1). The mixture was then transferred into a 3 mm diameter, 1.5 mm high, thin wall metal disc and cured for 10 minutes at 130 °C on a hot plate. Both sides of each disc were then ground on a 1200 grit paper until a thickness of 90–100 μm was obtained. A 400 mesh Cu grid was then stuck on one side of the disc and the opposite side was dimpled to about 40 μm using a 515 SBT dimple grinder with 6 μm diamond paste. Finally the dimpled side of the disc was thinned in a Technoorg-Linda ion beam miller. A problem with this technique was that the glue tended to be sputtered away preferentially during ion milling. To reduce the differential thinning, ion milling was conducted at low incident angles (∼10°). Moreover, the Cu grid was used to support powder samples weakened by glue removal from thinning, allowing further thinning to continue. The resulting specimens showed relatively large thin areas.
The variation of the hardness of powder particles with milling time was determined by microhardness measurements using a Vickers indenter (HV) at a load of 10 g and dwell time of 10 s. This produced a square impression with an average diagonal length of 7–3 μm, depending on milling time, which is at least twenty times smaller than the powder particle size. To eliminate errors caused by cold working of the surrounding area of indentations, the indentations were never made any closer than three times the diameter of an impression to an existing impression. 5–10 indentations were made on each sample to obtain an average value of microhardness.
Fig. 1 shows SEM micrographs of as-received Ni and Nb powders, together with Ni60Nb40 powder particles after different milling times.
Elemental Ni and Nb powder particles had a size distribution of ∼25 to 50 μm and ∼5 to 110 μm respectively. The average powder particle size increased to 280 μm after MA for 2.5 h, Fig. 1 (c). As shown in Fig. 1 (d), increasing milling times to 5 h led to a further increase in particle size. After 5 h of milling time the particles were varied in size, with an average of 430 μm, and also varied in shape. As illustrated in Fig. 1 (e–g), for milling times longer than 5 h the particle size decreased and their shape became progressively more uniform so that the final product after 85 h of milling, Fig. 1 (h) had a nearly spherical morphology and a narrow size distribution with a mean size of 45 μm. A higher magnification SEM micrograph of the particle surface, after 85 h of milling time, is shown in inset of Fig. 1 (h).
3.2 Structural changesFig. 2 shows XRD traces from Ni60Nb40 powder as-received and after different milling times.
XRD traces from Ni60Nb40 powder as-received and after different milling times.
The XRD traces of the as-received powders showed diffraction peaks from pure crystalline Nb and Ni. The sharp crystalline diffraction peaks broadened and their intensities decreased progressively during MA. As is evident in Fig. 2, after a milling time of about 20 h, a halo, which is attributed to the amorphous phase, developed on the XRD traces at a position between the crystalline Nb(110) and Ni(111) peaks. After a milling time of 30 h the high angle crystalline diffraction peaks disappeared into the background and only broadened Nb(110) and Ni(111) peaks along with an amorphous halo were apparent on the XRD traces. Further milling resulted in the crystalline Ni and Nb peaks gradually decreasing in size and finally vanishing. At the same time the amorphous halo gradually grew until the XRD traces appeared to show a completely amorphous structure without any indication of additional phases. The crystalline diffraction peaks did not shift during MA even at the longer times. Also XRD showed no intermediate crystalline phase as a precursor to the amorphous phase.
A problem with MA process, in particular for long processing time, is contamination. The contamination of powders during MA can arise from the milling media (Fe, Cr) and the atmosphere (O2, N2). The contamination by oxygen and nitrogen is of more importance during MA processing of reactive powders such as Zr, Nb, and Ti. The contamination of powder by oxygen and nitrogen in MA can be due to the oxygen and nitrogen absorbed on powder surfaces, residual oxygen and nitrogen in the bowel and oxygen and nitrogen picked up by leakage of air into the bowl during milling due to inadequate sealing of the bowl.
The XRD traces from as-milled samples did not show any indication of an additional phase (e.g. an oxide or nitride) as a result of reaction of powders by O2 or N2. Contamination by wear of milling media is not a serious problem during MA of metallic powders because of adhesion of powder particles on milling surfaces which limits their wear.
3.3 Microstructural evolutionsFig. 3 shows cross sectional SEM images of Ni60Nb40 powder particles after various milling times. As shown in Fig. 3 (a) the MA process first produced a layered microstructure consisting of cold welded Ni (dark area) and Nb (bright area) layers. After 2.5 h of milling time the thickness of the layers was quite non uniform from one particle to another and across the cross section with an average thickness of ∼11 μm. Increasing processing time to 5 h and then 10 h resulted in a finer and more uniform layered structure as shown in Fig. 3 (b) and (c). The average layer thickness after 5 h and 10 h of milling times were about 8 and 3 μm respectively. The microstructure after 15 h is of particular interest. As shown in Fig. 3 (d) the microstructure at this stage of MA consisted of two zones. Zone I was at the edges of the particles where no layered structure appeared within the resolution of the SEM. Zone II was in the internal part of the particles, with a clear and fine scale layered structure. The average layer thickness of ∼3 μm after 10 h of milling time was reduced to a value of ∼1.5 μm after 15 h of milling time. Most of the particles showed such a two zone microstructure, although in some particles the layered structure continued to be observed throughout the cross section. Milling for 25 h extended the size of zone I, however, the internal part still continued to exhibit a layered structure with a layer thickness which did not decrease further, and remained similar with that for 15 h. Fig. 3 (f) shows a higher magnification SEM of a particle after MA for 25 h. Traces of the layered structure were still evident in zone I as darker inhomogeneities. As MA proceeded further, the zone I gradually grew further inwards as shown in Fig. 3 (g), and the layer thickness in zone II did not refine significantly. Apart from a few particles the microstructure was completely featureless in the SEM after 85 h of milling time when XRD traces suggest a fully amorphous structure.
Cross-sectional SEM images of Ni60Nb40 powder particles after different milling times.
Fig. 4 shows the average layer thickness of particles centres versus milling time. The average layer thickness progressively decreased on continued milling and then reached a constant value after ∼20 h of milling time corresponding to the start time of amorphisation reaction. It was also noted that there is a tendency to form a finer microstructure at the edges of particles because of non-uniform deformation occurring in MA.
Variation of average layer thickness as a function of milling time for Ni60Nb40 powder particles.
To provide detail of microstructural evolution at the edges of particles during MA process TEM observations was used. Fig. 5 (a–c) shows bright field (BF) images and selected area diffraction patterns (SADPs), in the insets of figures, from individual Ni and Nb layers after MA for 5 h. Generally the dark (or white) areas on TEM images indicate the crystallineregions with the same orientations. A heavily deformed structure characterized by arrays of lattice defects (mainly dislocations) was observed for both Ni and Nb after 5 h of milling time. The original grain boundaries are still visible at this time of milling. The Nb grains in Fig. 5 (b) were long in one direction in contrast, Ni had equiaxed grains as shown in Fig. 5 (a). The corresponding selected area diffraction SADP for both Ni and Nb exhibited the Debye-Scherrer rings characteristic of a fine grain structure.
TEM images and corresponding selected area diffraction patterns of individual Ni and Nb layers after 5 h of milling time.
XRD peak broadening suggested grain sizes of 163 nm and 66 nm for Ni and Nb respectively after 5 h of milling time which are several order of magnitude smaller than the original Ni and Nb grain sizes. It appears that the original grains were split into the several sub-grains by arrays of dislocations after 5 h of milling time. This sub-grain structure can be also seen in Fig. 5 (c) in which the arrays of defects formed a cell structure within a large Nb grain. These small sub-grains, in fact, contribute towards the broadening of X-ray diffraction peaks and also contribute to the continuous Deby-Scherrer rings in the SADP. Grain boundaries within original grains were not well-defined after 5 h of milling time. This did not allow direct measurements of the Ni and Nb grain sizes on TEM micrographs.
A typical TEM image from edges of powder particles after 15 h of MA is shown in Fig. 6 (a). A fine layered structure with an average thickness of ∼60 nm was observed.
TEM images and corresponding selected area diffraction patterns of Ni60Nb40 powder particles edges after 15 h of milling time.
The corresponding SADP included the Debye-Scherrer rings of fcc (Ni) and bcc (Nb) structures without any indication of a third phase (e.g. an intermetallic compound or an amorphous phase) which is consistent with the XRD results. Arrays of dislocations within the layers were still present after 15 h of milling time however, the original grain boundaries disappeared. The sub-grain boundaries in layers were not well-developed after 15 h of milling, making measurements of the Ni and Nb grain sizes impossible. XRD peak broadening however, suggested a grain size of 31 nm for Ni and 18 nm for Nb after 15 h of milling time which are considerably smaller than the average layer thickness, ∼60 nm. Such a nanosized layered structure was found for most of the particles after 15 h of milling time, although occasionally regions with a coarser layered structure were also found at the edges of some particles for this sample.
A typical TEM image of these regions is presented in Fig. 6 (b). An extensive Ni layer with a well-defined nanocrystalline structure was observed in bright field (BF) and dark field (DF) images. The black areas on DF image indicate the grains with the same orientations. The dark field (DF) image using the crystalline Ni(111) diffraction ring, suggested that this area had an average grain size of ∼25 nm which accords with the average grain size 31 nm estimated using XRD peak broadening.
After 20 h of MA, the amorphisation reaction starts. As stated earlier, amorphous phase develops at the edge of particles and proceeds inwards as milling time increases. At this stage the particle structure was inhomogeneous with respect to the progress of the amorphous phase. For instance Fig. 7 shows an area of this sample with a greater fraction of amorphous phase at the interface between the crystalline Ni and Nb layers.
TEM images and corresponding selected area diffraction patterns of Ni60Nb40 powder particles edges after 20 h of milling time.
On continued milling the amorphisation reaction proceeded further. After 50 h of milling the microstructure at the edges of particles consisted of a large number of very small, unreacted Ni and Nb particles with a size of ∼3 nm embedded in an extensive amorphous matrix, Fig. 8. The structural inhomogeneity observed in the early stages of milling was much reduced after 50 h of milling.
TEM images and corresponding selected area diffraction patterns of Ni60Nb40 powder particles edges after 50 h of milling time.
Finally Fig. 9 shows the microstructure of a Ni60Nb40 powder sample after 85 h of milling. The uniform contrast of the BF and DF images along with a single diffuse ring on the SADP all suggested a fully amorphous structure at the edges of particles in agreement with the XRD results. This fully amorphous structure was found everywhere at the edges of all the particles.
TEM images and corresponding selected area diffraction patterns of Ni60Nb40 powder particles edges after 85 h of milling time.
The dependence of hardness on the microstructure of a material makes it a useful tool to study the microstructural changes occurring during mechanical alloying. Fig. 10 plots the average value of microhardness at the edges and centres of Ni60Nb40 powder particles after different milling times.
Microhardness value at the edges and centres of Ni60Nb40 powder particles after different milling times.
In the early stage of MA, up to 5 h, the microhardness values at the edges and centres of particles were identical. In contrast, the edges and centres of particles start to show different values of microhardness, 1160 and 814 Hv respectively, after 10 h of milling time. After 15 h of milling time, the value of microhardness at the particle edges, increased to 1280 Hv, and that for particle centres increased to 910 Hv. After 25 h of milling time the microhardness value of the particle edges increased even further, to 1670 Hv, and that for the particle centres also increased further to 1070 Hv. This remarkable increase in microhardness at the edges is consistent with the presence of an amorphous phase, as detected by TEM. The large difference in microhardness between the edges and centres of the particles was also observed after 50 h of milling time. At this time the microhardness value of the particle edges had increased to 1820 Hv as a result of the continued gradual amorphisation process while that for particle centres increased only slightly to 1080 Hv. After 85 h of milling, when XRD and TEM suggested a fully amorphous structure, the edges and centres of the particles both had the same, very high microhardness value of 1960 Hv.
Changes in particles morphology and microstructure during MA of ductile metal powders are produced by two simultaneous processes; cold welding and fracturing.
During mechanical alloying clusters of particles are trapped between colliding balls and undergo a high level of impact. If the impact stresses are sufficient, the powder particles plastically deform and flatten. As the powder particles are pressed together their surface area increases and the surface oxide films rupture, consequently exposing clean underlying metal. When these fresh surfaces of particles come in contact a metal bond is formed.
After a period of milling, particles deform to the extent that cracks initiate, propagate and ultimately fracture the particles.
The extent of these two events is determined by the mechanical properties of the elemental powders, such as ductility, yield stress and hardness, as well as the magnitude of the impact provided by colliding balls.
Three stages were observed in analysing the powder particle size during MA of Ni60Nb40 and powder mixture. The initial particle size first increased reaching a maximum after 5 h of milling time. In the second stage the particles rapidly decreased in size. This was followed by a third stage; steady-state stage for milling times longer than ∼20 h in which the particle size remained almost constant. These three stages and associated microstructural changes can be analysed by considering the relative rates with which the cold welding and fracturing processes occur, as shown schematically in Fig. 11.
Schematic illustration of rate of cold welding and fracturing of powder particles during mechanical alloying process.
In first stage (0–5 h) the cold welding process dominates and as a result the powder particle size continuously increases. This stage can be termed the agglomeration stage. The flattening and cold welding of powder particles during this period of MA result in the development of a layered structure consisting of Ni and Nb layers. The cross sectional SEM images showed that in the first few hours of MA the layers are coarse and vary in thickness over the particle cross sections. As MA proceeds further, the layer thickness progressively refines and becomes more uniform as a result of the repeated cold welding and fracturing of particles as shown in Fig. 3 and Fig. 4. The ultrafine layered microstructure developed during ball milling provides extensive interfaces suitable for any potential reaction (e.g. amorphisation, formation of crystalline compounds, ...) between constituents at longer milling times or during subsequent processing (e.g. hot press, hot extrusion, thermal spray).
Microhardness measurements on Ni60Nb40 particles showed a progressive increase in hardness value during the first stage of MA as a consequence of work hardening. This leads to a decrease in the ductility of powder particles and therefore, an increased tendency for particle fracture. Thereby, the first stage is followed by the fragmentation stage which lasts from ∼5–20 h of milling time. During the fragmentation stage the fracturing of particles occurs more readily than cold welding and as a result the powder particle size decreases. The layered structure also refines further in the second stage. TEM images from Ni60Nb40 particles after 15 h of milling time revealed a nanoscale size layered structure, with a typical thickness of ∼60 nm. The faster refinement of the layered structure at the edges of particles implies that the edges are subjected to greater plastic deformation than the centres, consistent with high deformation rates in MA. The microhardness measurements provided further support for this non-uniform plastic deformation. After 15 h of milling time the edges of the particles attain a higher hardness (1280 Hv) compared with the centres (910 Hv) because of higher work hardening. As will be discussed next section this inhomogeneity in plastic deformation results in the edges amorphising first, before the centres.
After around 20 h of milling time there is a steady-state stage in which the powder particle size does not change significantly. Interestingly the commencement of the steady-state stage, after ∼20 h of milling time, coincides with the development of an amorphous phase. The amorphisation reaction was observed to start at the edges of particles, where a nanoscale size layered structure is first formed, and continues inwards. It is believed that during the steady-state stage there is a balance between the frequencies of cold welding and fracturing processes so that the average particle size remains unchanged (Dunlap et al., 2000). This mechanism of continuous equal amount of the fragmentation and cold welding of particles in the steady-state stage should lead to a random distribution of the amorphous phase throughout the whole of the particle volumes. The observed progress of the amorphisation reaction from the edges towards the centre of particles in the present work however, suggests that no significant fracturing or cold welding occur during the steady-state period. The start time of the steady-state stage and the formation of an amorphous phase at the edges of particles appeared to be similar, after ∼20 h of milling time, indicating that the amorphous shell at the edges of powder particles prevents further fragmentation. Amorphous alloys are extremely hard, but also can deform substantially in compression, therefore the amorphous particle edges provide a strong and resilient layer preventing fracture during milling. The earlier mechanism based on the balance between the cold welding and fracturing processes could be encountered in high energy mills where a much higher force is imposed on particles and in other cases where a resilient amorphous outer layer is not formed.
MA of Ni60Nb40 alloy in low energy centrifugal ball mill consists of three stages; agglomeration, fragmentation and steady-state. During the first stage of agglomeration successive cold welding of powder particles leads to a continuous increase in particle size. In this stage a layered structure is formed, with a progressively refined layer thickness with increasing milling time. The agglomeration stage is followed by the fragmentation stage as work hardening limits ductility and further cold welding. During fragmentation stage fracturing of particles dominates, leading to a rapid decrease in particle size as well as a continuing refinement of layered structure. The grain size, microstructure and microhardness value at edges and centres of particles were different at early stage of MA which is consistent with the non-uniform deformation occurring in MA. This causes the particle edges to amorphize prior to the centres. The amorphous phase at the edges of particles gradually proceeds inwards with increasing milling time until the whole particles became uniform with respect to the microstructure and microhardness value i.e. the edges and centres of the particles both had the same features. However, the amorphous shell provides a strong and resilient layer, preventing further fracture. Thereby, the fragmentation stage is followed by a third stage; steady-state stage in which powder particle size remains constant. Even though the impact force is not high enough to rupture the amorphous layer at the edges, the regions underneath can be plastically deformed and then transform to the amorphous phase. Continued milling progressively refines the grain size and increases the internal strain introduced by increasing density of crystal defects. An ultrafine layered structure with a typical thickness of 30 nm, containing nanoscale size grains with a typical size of 15 nm and a high density of dislocations, develops prior to amorphization reaction. This observation suggests that numerous high speed diffusion paths are necessary to allow a high diffusion rate at low temperature and therefore permits the amorphization reaction to take place kinetically.
Mohammad Hossein Enayati
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M. H. Enayati (PhD, CEng, MIMMM) is a full Professor of materials science at Isfahan University of Technology. Prof. Enayati holds a bachelor’s degree in metallurgical engineering from the Isfahan University of Technology,-Iran, a master’s degree in materials engineering from Shiraz University-Iran, and a doctorate in materials science from the University of Oxford-UK. He is the former Dean of Office for Research Affairs, Graduate Program Advisor, Deputy of Research Affairs, member of editorial board of Journal of Advanced Materials in Engineering and founder director of Nanocenter at Isfahan University of Technology. Prof. Enayati’s research focuses on the nanostructured and amorphous materials, mechanical alloying and synthesis of advanced materials for thermal spray coating. He has authored several books and numerous articles, is a frequent presenter at professional conferences, has served in editorial capacities for several leading journals in his field, and has received research funding from a wide variety of agencies. His research has been recognized with numerous awards, most recently from the Iran Nanotechnology Initiative Council and Iranian Nanotechnology Society.
