The influences of La doping method on the microstructure and mechanical properties of Mo alloys

Int. Journal of Refractory Metals and Hard Materials 51 (2015) 301–308

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The influences of La doping method on the microstructure and mechanical properties of Mo alloys Xiaoqing Yang a,b, Hua Tan a, Nan Lin a, Zhixiang Li b,c, Yuehui He a,⁎ a b c

State Key Laboratory of Powder Metallurgy, Central South University, Changsha, China Zigong Cemented Carbide Co., Ltd, Zigong, China School of Material Science and Engineering, Central South University, Changsha, China

a r t i c l e

i n f o

Article history: Received 18 September 2014 Received in revised form 22 April 2015 Accepted 28 April 2015 Available online 29 April 2015 Keywords: Doping method Molybdenum alloy La2O3 Microstructure Mechanical properties

a b s t r a c t In this paper, the effects of the La doping method on the microstructure and mechanical properties of Mo alloys were investigated. The powders were fabricated by the industrial produced pure Mo powder, solid–solid doping method (S–S), solid–liquid doping method (S–L) and liquid–liquid doping method (L–L), respectively. In addition, the powders were processed into sintered Mo, rotary swaged Mo and Mo wire. The results indicated that the powders became finer and the grain boundaries of the Mo prepared by S–L(S–L Mo) and L–L (L–L Mo) had the ability to migrate and wrap the La2O3 particles into the grain interior during sintering with the addition of La. The tensile strength and elongation of the rotary swaged L–L Mo were the highest and the tensile elongation was raised to 42% at room temperature. Moreover, the L–L Mo wire had the highest tensile strength after the treatment at all the annealing temperatures and the S–L and L–L Mo wire possessed higher recrystallization temperature than the Mo prepared by S–S (S–S Mo) and pure Mo wire obtained. © 2015 Elsevier Ltd. All rights reserved.

1. Introduction Molybdenum (Mo) is a refractory metal, having low coefficient of linear expansion, excellent thermal conductivity, satisfactory electrical conductivity and high elastic modulus, and is thus very popular in metallurgy, oil, electronic product and nuclear industry. In industry, molybdenum wire is very important in the production of high temperature electric and line cut machining. The production of pure molybdenum wire has been very mature and the performance of all aspects has been satisfactory [1–3]. However, in order to improve the property, adding the dopants is an effective method, for example, Si, Al and K are satisfactory dopants for molybdenum wire to improve the high temperature mechanical properties and recrystallization temperature. The AKS-doped molybdenum wire contains small volume fractions of aluminosilicate and potassium aluminosilicate dopant particles in addition to potassium bubbles and tubes/ellipsoids. Rare earth element is another choice. Lanthana-doped molybdenum has been reported to have recrystallization temperatures above those obtain through AKS doping. Lanthana particles are stable within molybdenum to elevated temperatures because lanthanum has very limited solubility in molybdenum. Besides the doping elements, method of doping is another vital factor. In general, the doping methods can be broadly classified as solid–solid (S–S), solid–liquid (S–L) and liquid–liquid (L–L) doping methods. Prior studies on the doping method have shown that the L–L doping is ⁎ Corresponding author. E-mail address: [email protected] (Y. He).

http://dx.doi.org/10.1016/j.ijrmhm.2015.04.034 0263-4368/© 2015 Elsevier Ltd. All rights reserved.

the most uniform method which has excellent performance because it makes the mixed powder forming the core–shell structure so that the La2O3 particles can be distributed in the grain interior. While the S–L doping is better than S–S doping, due to the dispersion effect of the liquid can refine the La2O3 particles [1,4–10]. This study focuses on the lanthanum doped molybdenum alloy and the manufacture processing from powder to wire. Though the molybdenum wire doped with lanthanum has been industrialized, most factories use solid–liquid doping method, even solid–solid doping method which just could meet the basic requirements. However, it is necessary to study the liquid–liquid doping method to improve the comprehensive operational performance. This study began from powder process, sintered Mo, rotary swaged Mo to Mo wire, compared to the differences of the microstructure and mechanical properties between different ways of doping and explained the mechanism. It is helpful and significant for future industrial production. 2. Experimental procedures The MoO2 powder for the experiment was reduced from the ammonium molybdate via one reduction, and the pure Mo powder was reduced from MoO2 powder. The S–S powder was fabricated by the S–S doping method that was putting the MoO2 powder blended with the La2O3 powder (2.1 μm, FSSS), then through reduction. The S–L powder was fabricated by the S–L method that was putting the MoO2 powder blended with the lanthanum nitrate solution, then through agitation, heating, evaporation and reduction. Homogeneously, the L–L powder

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was fabricated by the L–L method that was putting the ammonium molybdate solution blended with the lanthanum nitrate solution and the ammonia, then through agitation, heating, evaporation and reduction. The mass fraction of La is 0.3% for all the mixed powders. The impurity content (mass fraction) of the ammonium molybdate, lanthanum nitrate and the La2O3 powder is given in Table 1. All the powders were statically cold pressed into cylindrical compact (20 mm in diameter), and sintered at 1960 °C for 6 h in flowing dry hydrogen. The impurity content (mass fraction) of the sintered bars is given in Table 2; the diameter and the density of the sintered Mo bars are given in Table 3. The cylindrical compact was thermomechanically processed into a rod by rotary swaging from the sintered diameter to 15.2–14.0–12.0– 11.0–10.0–9.0–8.0 mm, and the preheat temperature for the rotary swaging was 1250 °C. Finally, the rod was drawn into the molybdenum alloy wire with the diameter of 0.68 mm at 650–700 °C. The wires were carefully heat treated at 900–1700 °C (every 100 °C) in a dry hydrogen atmosphere for 0.5 h. The morphologies of the powders were observed by a scanning electron microscope (SEM, JSM-6360LV, 20 kV), a laser light scattering analyzer (Mastersizer 2000) and a Fisher sub-sieve sizer. The density of the sintered bars was measured by the Archimedes method. The microstructure of the sintered Mo and rotary swaged Mo rods was carried out by the SEM of polished specimens etched with Murakami's etch (an aqueous solution of potassium ferricyanide and sodium hydroxide) and a transmission electron microscope (TEM, JEM-2100, 200 kV) of the foils conducted by ion beam thinner. The grain size of the sintered Mo and rotary swaged Mo was conducted by a linear intercept method on metallographic samples. The Mo wires of processing state were observed by SEM of polished specimens. Room temperature tensile tests of the rotary swaged Mo rods (as-swaged condition) and Mo wires of all annealing temperatures were performed using an Instron 3369 testing machine at a constant strain rate of 1.5 mm/min. Results of tensile strengths and elongation to failure were recorded and calculated by computer. 3. Result and discussion 3.1. Morphologies and properties of the powder The morphologies of different powders are given in Fig. 1. These figures show that the S–L powder has the finer size than the S–S powder

Impurity content

Pure Mo

S–S Mo

S–L Mo

L–L Mo

La (ω. %) K (ω. ppm) Fe (ω. ppm) O (ω. ppm)

NMa 25 39 10

0.34 27 41 600

0.28 47 44 600

0.30 51 38 600

a

Not measured.

and the L–L powder size is the finest. The particle size distribution is shown in Fig. 2; the characteristic parameters of size distribution and the FSSS number are shown in Table 4. From the distribution, the characteristic parameters and the FSSS number, the L–L powder possesses the finest size and the largest specific surface area which in line with the morphologies observed from the figures. In addition, the L–L powder has narrow particle size distribution from 1–7 μm while the distribution of the pure Mo powder is from 1–40 μm. According to the chemical vapor transport (CVT) model, the decomposition of the educt is followed by formation of an intermediate gaseous transport phase. This transport phase is deposited on a nucleus of the product. This way, the grain morphology of the product is generated completely new. The grain size distribution of the product is determined by the conditions of deposition (heterogenous/homogenous nucleation, nucleus growth) from the gas phase [11]. The La2O3 particles can provide more nucleus to form more fine particles and inhibit grain growth during reduction. However, in the S–L doping, particles of lanthanum nitrate are coated on the surface of the Mo powder after evaporative crystallization. Then the particles of lanthanum nitrate are decomposed to La2O3 particles during reduction. And the La2O3 particles are finer and provide more nucleus than the S–S doping did. In the L–L case, the process ensures mixing on the molecular level, and encourages heterogeneous nucleation to form core–shell structures of very fine lanthanum-containing particles in the ammonium dimolybdate crystals. The outcome is to form the mixed powder with very fine La2O3 particles inside the Mo powder after reduction. In addition, due to the finer La2O3 particles, more nucleus is provided to form finer Mo powder [6]. Consequently, the S–L powder has the finer size and the larger specific surface area than the S–S powder, and the L–L powder has the finest size and largest specific surface area. 3.2. Microstructure of the sintered molybdenum

Table 1 The impurity content (mass fraction) of the starting powder. Element

Ammonium molybdate, b/= ppm

Lanthanum nitrate, b/= ppm

Oxide

Lanthanum oxide, %

Fe Al Si Mn Mg Ni Ti V Co Pb Bi Sn Cd Sb Cu Ca Cr P K As Na Chloride Sulfate

5 1 5 1 1 1 5 5 1 1 1 1 1 10 1 5 5 5 10 10 1 NMa NMa

5 5 NMa 5 5 NMa NMa NMa NMa 5 NMa NMa 5 NMa 5 5 NMa NMa 10 NMa 5 50 50

CeO2 Pr6O11 Nd2O3 Sm2O3 Y2O3 Fe2O3 SiO2 CaO CuO MnO2 NiO PbO CoO Cr2O3

0.0003 0.0003 0.0002 0.0002 0.0002 0.0001 0.0025 0.0010 0.0001 0.0001 0.0002 0.0002 0.0001 0.0001

a

Table 2 The impurity content (mass fraction) of the sintered bars.

Not measured for no detection method due to limited experiment condition.

Fig. 3 shows the SEM microstructure of the sintered molybdenum. It can be seen from the figures that the S–S Mo possesses finer grain size than the pure Mo and the grain size of the S–L Mo and L–L Mo is bigger than the S–S Mo. The figure also shows the second phase particles; Fig. 3 (e), (f) and (g) shows the EDS analysis of the second phase particles in the sintered molybdenum. From the EDS, Oxygen and Lanthanum are the main components in these second phase particles. And most of the second phase particles are distributed on the grain boundaries in the S–S Mo and are distributed in the grain interior in the S–L Mo and L–L Mo. Moreover, the second phase particles on the grain boundaries are coarser than those in the grain interior. For the S–S Mo, most of the La2O3 particles are coarse particles (discussed above) that have enough ability to hinder the migration of the grain boundaries during sintering. As a result, most of the La2O3 particles in the S–S Mo are coarse and distributed on the grain boundaries. While for the S–L Mo, most of the La2O3 particles are fine that can be wrapped into the grain interior for Table 3 The diameter and density of the sintered bars.

Diameter (mm) Density (g/cm3)

Pure Mo

S–S Mo

S–L Mo

L–L Mo

15.82 9.97

15.64 9.93

16.96 9.63

16.60 9.84

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Fig. 1. The morphologies of the powders. (a) pure Mo, (b) S–S Mo, (c) S–L Mo and (d) L–L Mo.

the migration of the grain boundaries during sintering. In addition, the powder of large specific surface area has high sintering activity that indicates the high grain boundary migration rate and ability [12]. As discussed earlier, the S–L powder has fine particles and larger specific surface that give rise to the grain boundary migration rate and ability. As a result, most of the La2O3 particles in the S–L Mo are fine and distributed in the grain interior, and the grain size is larger than the pure Mo and S–S Mo. At last, for the L–L Mo, it contains the finest La2O3 particles and some mixed powders of core–shell structures that all contribute to the fine La2O3 particles and the distribution of the grain interior in the sintered samples. Consequently, the L–L Mo has the finest La2O3 particles and most of the La2O3 particles are distributed in the grain interior.

35

pure Mo S- S Mo S- L Mo L- L Mo

30

Cum%

25 20 15

3.3. Microstructure and mechanical properties of the rotary swaged Mo The transverse section and longitudinal section microstructure of the rotary swaged Mo can be seen from Fig. 4. The grains of the pure Mo and the S–S Mo are regular and uniform and the grain size is coarse after rotary swaging. The grains of the S–L and the L–L Mo are elongated and its size is fine and uneven after rotary swaging. Due to the fact that measuring the grain size in rotary swaged Mo is difficult because of the structure's complexity, and the difficulty of discriminating between grain boundaries and subgrain boundaries in etched samples, we can just estimate the grain size of the samples. Why does the grain shape and size have such big difference? It is because the preheat temperature for the rotary swaging is 1250 °C. The grains of the pure Mo and S–S Mo will be recrystallized at 1300 °C and 1500 °C, respectively (see the following analyses), while for the rotary swaged Mo, the recrystallization may take place for most of the grains of the rotary swaged pure Mo and for a great part of grains of the rotary swaged S–S Mo. So their grains are regular and uniform. However, for the S–L and L–L Mo, their recrystallization temperature is about 1800 °C, so the grains will not recrystallize. Moreover, the deformation of the rotary swaging leads to the dislocation multiplication and movement. As a result, the subgrain structure is formed. The Hugo R.Z. Sandim' study [13] also has this

10 Table 4 The characteristic parameters of size distribution.

5 0 0.1

1

10

Size(µm) Fig. 2. The particle size distribution curve of the powder.

100

Pure Mo S–S Mo S–L Mo L–L Mo

D50 (μm)

D(4,3) (μm)

D(3,2) (μm)

S.S.A. (sq.m/c.c.)

FSSS (μm)

11.89 9.45 4.55 3.09

12.36 9.94 4.70 3.10

8.90 6.78 3.49 2.66

0.67 0.89 1.72 2.26

3.7 3.5 3.0 2.6

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Fig. 3. The SEM microstructure of the sintered molybdenum. (a) pure Mo, (b) S–S Mo, (c) S–L Mo, (d) L–L Mo; (e), (f) and (g) EDS of the second phase particles for (b), (c), and (d).

structure. It also can be seen from the figure that the pores of the rotary swaged Mo become fewer and smaller than the sintered Mo. Due to the deformation of the rotary swaging, the pores within the material can be shrank and welded together. The tensile strength and elongation to failure of the rotary swaged Mo are shown in Fig. 5. The room temperature tensile strength of the pure Mo is 562.8 MPa and the elongation is 22.4%. The rotary swaged Mo doped with La has higher strength and elongation. Among them, the L–L Mo has the highest tensile strength of 637.3 MPa and the highest elongation of 42.0%. For the reason that when the deformation begins from one grain to the other, it is difficult to cross the grain boundaries which possess complex structure, and this process of deformation needs large energy. Consequently, the finer grain size causes the more grain boundaries and the more grain boundaries result in the higher strength and elongation. Therefore, the strength of the S–L and the L–L Mo was higher than the S–S Mo and the strength of the S–S Mo was higher than the pure Mo. This is one reason for the higher strength and elongation of the S–L and L–L Mo than the S–S Mo. And another vital reason is that in the S–S Mo most of the La2O3 particles distribute in the grain boundaries, while in the S–L and L–L Mo most of the La2O3 particles distribute in the grain interior. The TEM microstructure of the La2O3 particles in the L–L Mo is shown in Fig. 6. From Fig. 6(a) the La2O3 particles are distributed in the grain interior and about 100–200 nm. Fig. 6(c) and (d) shows the EDS and TEM diffraction spot, respectively. Based on previous studies, the intragranular particles can generate, pin down and thus accumulate dislocation within the grain [6]. For the S–L and L–L Mo, the La2O3 particles that distribute in the grain interior have the dual effect of blocking dislocation and storing dislocation during deformation. It can be seen from Fig. 6(b) that there are a few dislocations tangled around the La2O3 particles. During the deformation, the dislocation constantly generated and moved, when it

crossed the grain, the La2O3 particles could block and pin down the dislocation, and then accumulate the dislocation within the grain. The dislocations tangled around the La2O3 particles could increase the resistance to slippage and give rise to the work hardening and toughness increasing. In order to compare the strengthening mechanism of the La2O3 particle dispersion and the grain refinement, we reference Zhang's model to analyze the two strengthening mechanisms. In Zhang's model the yield stress can be written: σ y ¼ σ M þ σ P þ k=D1=2 where σy, σM and σP are the yield stress, the matrix strength and the strength due to the fine particles. k/D1/2 is the grain size strengthening term, where k is the Hall–Petch slop and D is the grain size. According to the Orowan–Ashby equation and Zhang's derivation, the increase in yield strength due to the particles is presented as follows: σ P ¼ σ OR ¼

  mμb Φ rffiffiffiffiffiffi ㏑ 2b π −1 ð1:18Þ  2π  Φ 6f

where m is the Taylor factor, μ is the shear modulus, b is the Burgers vector, Φ is the particle size and f is the volume fraction of the oxide lanthanum particles that can pin down the dislocations. So the yield stress can be written as follows:

σy ¼ σm þ

  mμb Φ k þ pffiffiffiffi : rffiffiffiffiffiffi ㏑ 2b π D −1 ð1:18Þ  2π  Φ 6f

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Fig. 4. The transverse section of (a) pure Mo, (b) S–S Mo, (c) S–L Mo and (d) L–L Mo, and longitudinal section microstructure of the rotary swaged Mo. (e) Pure Mo, (f) S–S Mo, (g) S–L Mo and (h) L–L Mo.

It can be concluded that the yield stress increases with the grain size decreasing and the volume fraction of the oxide lanthanum particles increases from the equation. The Mo alloys doped with La have higher strength than the pure Mo for the two strengthening mechanisms. In addition, the strength of the S–L and L–L Mo is higher than the S–S Mo for their finer grain size and higher volume fraction of the La2O3 particles. Moreover, the strength of the L–L Mo is higher than the S–L Mo,

for the reason that the La2O3 particles are finer and more uniform than the S–L Mo and have more ability to hinder and pin down the dislocations. The intragranular La2O3 particles not only can increase the strength but also have positive effect to the elongation since the particles can help with the generation and storage of dislocations without initiating cracks localized at grain boundaries [6]. As a result, the S–L and L–L Mo have both higher strength and elongation than the S–S Mo.

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Fig. 5. The tensile strength and elongation to failure of the rotary swaged Mo at room temperature.

3.4. Microstructure and mechanical properties of the Mo wire The longitudinal section and cross section microstructure of the Mo wire are shown in Fig. 7. As shown in Fig. 7(a), (b), (c) and (d), the microstructure of the longitudinal section is the fiber grain along the drawing direction. As can be seen from the width of the fiber grain, the L–L Mo wire possessed the finest grain which could be observed from the cross section microstructure obviously shown in Fig. 7(e), (f), (g) and (h). The rotary swaged microstructure will be drawn out in the severe plastic deformation of the drawing process. Due to the fine grain size and the subgrain structure of the rotary swaged Mo, the L–L Mo wire has the finest fiber grain under the same deformation. Fig. 8 shows the tensile strength and elongation of Mo wires after annealing at 900–1700 °C. It can be seen from Fig. 8(a) that the tensile

a

c

strength decreases with the annealing temperature increase. For the L–L Mo wire, the tensile strength of the processing state was 1611 MPa. However, after annealing at 900 °C, the strength decreases to 1455 MPa and then becomes lower after annealing at the following temperature. The reason is that the annealing process can eliminate internal stress that exists in the material and at higher annealing temperature, the recrystallization may take place and change the grain from fiber to interlocked elongated grains [9], thus leading to the decrease of strength. From Fig. 8(a), after heat treatment at 900 °C, the tensile strengths of the pure Mo, S–S Mo, S–L Mo and L–L Mo are 1149 MPa, 1245 MPa, 1398 MPa and 1455 MPa, respectively. The change rule of the tensile strength of the Mo wire with the doping method is identical to the rotary swaged Mo. From the observation in the microstructure of the Mo wire, the grain size is an important reason as discussed earlier. The other reason is also the dispersion of the La2O3 particles. For the S–S Mo, the coarse La2O3 particles were elongated during drawing and broken up into the row of spherical particles during annealing [9]. For the S–L and L–L Mo, we speculated that only few particles of coarse La2O3 particles would be broken up, while the other fine particles would not experience this effect. Nevertheless, the L–L Mo wire still possessed the finest and most La2O3 particles within the grain. During the tensile deformation of the Mo wire, the La2O3 particles could still pin down dislocation and thus accumulate dislocation, consequently resulting in the high strength and elongation. However, the Mo wire possessed large residual stress after the drawing, and it revealed low elongation before annealing. The elongation of the Mo wire after annealing at 900 °C to 1700 °C is shown in Fig. 8(b). It can be seen from the figure that elongation increases with the annealing temperature increase at 900–1300 °C. After heat treatment at 1300 °C, the elongation of pure Mo gets to 24% and then decreases with the annealing temperature increase. And after annealing at 1500 °C, the elongation of the S–S Mo gets to 21.5% and then decreases. For the S–L and the L–L Mo, there is no highest elongation. According to the previous studies of recrystallization [5,9], the reason for the highest elongation is that the material recrystallized at the temperature of the highest elongation. It can be concluded that the addition of La can

b

d

062

136 134

Fig. 6. (a), (b) The TEM microstructure of the rotary swaged L–L Mo; (c) EDS of the second phase particle in (b); (d) the diffraction spot of the second phase particle in (b).

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Fig. 7. The longitudinal section microstructure of the Mo wires, (a) pure Mo, (b) S–S Mo, (c) S–L Mo and (d) L–L Mo; the cross section microstructure of the Mo wires, (e) pure Mo, (f) S–S Mo, (g) S–L Mo and (h) L–L Mo.

increase the recrystallization temperature. For the metal Mo of high stacking fault energy, the nucleation during the recrystallization is main through the climbing and gliding of the dislocation existed on the boundaries of the adjacent subgrain [9]. And the La2O3 particles dispersed in the grain boundaries and the grain interior could

block the motion of the dislocation. According to the “Johnson– Mehl equation”  π  Xr ¼ 1− exp − NG3 τ3 3

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4. Conclusions (1) The addition of the La can refine the powder, and the S–L powder has the finer size than the S–S powder, the size of L–L powder is the finest. (2) For the sintered molybdenum, the S–S Mo has the finest grain size, while the grain size of the S–L and L–L Mo is much bigger than pure Mo. In the S–L and L–L Mo most of the La2O3 particles are distributed in the grain interior and in the S–S Mo most of the La2O3 particles are distributed on the grain boundary. (3) For the rotary swaged molybdenum, the L–L Mo has the highest tensile strength and the highest elongation. (4) For the molybdenum wires, the tensile strength decreases with the annealing temperature increase. In addition, the L–L Mo has the highest strength and the S–L and L–L Mo have the highest recrystallization temperature.

References

Fig. 8. (a) The room temperature tensile strength and (b) elongation to failure of the molybdenum wires after annealing at 900–1700 °C.

where Xr is the recrystallized volume fraction, N the nucleation rate, G the growth rate and τ the annealing time [14]. The La2O3 particles can decrease N and G, thus decrease the recrystallization rate. And finer La 2 O3 particles will reduce the recrystallization rate more. Moreover, the coarse grain would lower N so that the recrystallization rate will decrease. As a result, the S–L and L–L Mo have the higher recrystallization temperature than the S–S Mo and the S–S Mo has the higher recrystallization temperature than the pure Mo.

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