- Email: [email protected]
Study on hydrogenation behaviors of a Mg-13Y alloy
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Available online at www.sciencedirect.com
ScienceDirect journal homepage: www.elsevier.com/locate/he
Study on hydrogenation behaviors of a Mg-13Y alloy Xiaoying Shi, Jianxin Zou, Chuan Liu, Lifang Cheng, Dejiang Li, Xiaoqin Zeng*, Wenjiang Ding The State Key Laboratory of Metal Matrix Composites, National Engineering Research Center of Light Alloy Net Forming, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, PR China
article info
abstract
Article history:
The Mg-13Y bulk alloy was prepared by conventional casting process and the Mg-13Y
Received 5 December 2013
powder was processed by ball milling using the casting alloy under the protection of
Received in revised form
argon. The hydrogenation thermodynamics, hydrogenation process and phase transitions
7 March 2014
were carefully investigated in the Mg-13Y powder alloy. It is shown that the Mg-13Y casting
Accepted 17 March 2014
alloy consists of Mg24Y5 phase and a-Mg containing yttrium which have different hydro-
Available online 22 April 2014
genation enthalpies, 195 kJ/mol H2 (by calculation) and 42 kJ/mol H2 (by Pressure eCompositioneTemperature experiment), respectively. The structure evolution and phase
Keywords:
transition in the Mg-13Y bulk alloy treated at 673 K and at 4 MPa for 40 h were observed by
Hydrogenation behavior
an optical microscopy (OM), a scanning electron microscopy (SEM), a transmission electron
Mg-13Y bulk alloy
microscopy (TEM) and X-ray diffraction (XRD). The large Mg24Y5 phase in the bulk Mg-13Y
Hydrogenation enthalpy
alloy could be destroyed into fine cuboid-shaped YH2 phases during the hydrogenation
YH2 cuboid phase
process, which is probably responsible for the improvement of mechanical properties of Mg-13Y alloy. Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Magnesium alloys, the lightest structural alloys developed so far, have great potential for lightweight applications, ranging from the portable electronic devices to automobile parts. However, the low strength limited the range of its application seriously. Therefore, lots of heat treatments, such as solution treatment [1] and ageing treatment [2], have been developed to improve its strength. Hydrogenation heat treatment applied to magnesium alloys is one of the thermo-chemical
treatments [3] which are not widely used. This is similar to nitriding [4] or carburizing [5] heat treatments implemented in steel and could be applied to thin-wall castings because hydrogen diffusion speed in bulk alloys is low [3]. Hydrogenation heat treatment was well applied to strengthen the MgeRE (Rare Earth)eZn casting alloys since the insoluble MgeREeZn phases at grain boundaries could be transformed into rare earth hydride and Zn atoms could dissolve into Mg matrix to enhance the strength of matrix. The newly formed rare earth hydrides exhibiting tiny particle-like morphology were thermodynamic stable and hard [6,7]. Large
* Corresponding author. Tel.: þ86 21 54742301; fax: þ86 21 34203730. E-mail addresses: [email protected] (X. Shi), [email protected] (J. Zou), [email protected] (C. Liu), chenglifang@china. com.cn (L. Cheng), [email protected] (D. Li), [email protected] (X. Zeng), [email protected] (W. Ding). http://dx.doi.org/10.1016/j.ijhydene.2014.03.115 0360-3199/Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
8304
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
improvements of ultimate tensile strength (UTS) (from 160 MPa to 292 MPa) and elongation (from 2.0% to 12.1%) compared to the as-cast alloy were observed in the ZM8 (Mg6 wt%Zn-2.5 wt%RE-0.5 wt%Zr) alloy treated at 753 K (480 C) and 1 atm pressure of hydrogen for 24 h. In addition, the method of hydrogen heat treatment followed with ageing treatment was more effective at strengthening this alloy. The UTS of aged alloy reached 316 MPa and the tensile yield strength (TYS) of aged alloy increased to 223 MPa from 129 MPa in the as-cast alloy [7]. Another important issue about hydrogenation which can not be ignored is that the method of hydrogenationedisproportionationedehydrogenationerecombination (HDDR) [8e12] could reduce the grain size of MgeAl alloys. In previous work, Takamura et al. [13e15] investigated grain size refinement in MgeAl-based alloy powders by hydrogenation treatment and found that the grain size of AZ31 alloy powders was reduced from 50 to 300 mm to about 100 nm in the case of the heat treatment at 623 K (350 C) under a hydrogen pressure of 7 MPa for 24 h. In addition, AZ31 alloy plate was also conducted HDDR process at 723 K (450 C) and its grain size was reduced to less than 500 nm, while the layer thickness being treated was limited to the range of about 20 mm from the surface after 36 h at 7 MPa pressure of H2 [13]. Mg matrix was hydrogenated into MgH2 in the HD process and then the MgH2 was dehydrogenated into Mg in the DR process. Microstructure and strengthening mechanism of hydrogenated MgeRE alloys were not investigated carefully and completely. It is significant to understand the microstructure evolution, phase transformation and corresponding thermodynamics during the hydrogenation process in order to make good use of this method and design technological parameter. In the present work, Mg-13Y binary alloy was selected as an example to investigate the microstructure evolution and strengthening mechanism of MgeRE alloys heat-treated in the hydrogen. Since the diffusion speed of hydrogen in the bulk alloy is very low, the hydrogenation process was firstly investigated in the Mg-13Y powder alloy. According to the Refs. [7] and [13], the hydrogenation behaviors of MgeRE based bulk alloy and MgeAl based bulk alloy are totally different. Therefore, the hydrogenation behaviors of Mg-13Y bulk alloy and Mg-8Al bulk alloy were also directly compared in this paper.
Experimental procedures Mg-13Y (Mg-12.85 wt% Y) and Mg-8Al (Mg-7.68 wt% Al) alloys were melted in steel crucible under a cover gas mixture of SF6 and CO2 using pure Mg, Mg-25Y (wt%) master alloy and pure Al. The melt was poured at 1003 K (730 C) into the mould preheated at 473 K (200 C). Bulk specimens with dimensions of 30 mm 10 mm 2 mm were cut from ingots by wire electrical discharge machining (WEDM) and removed oxide scale from the surface by using SiC paper in the Ar-filled glove box. The coarse powder samples used for thermodynamics researches were prepared in air by punching machine but quickly transferred into the Ar-filled glove box in approximately one minute. Under argon protection, the fine powder
samples were produced from coarse powder using a QM-3SP2 planetary ball milling machine made by Nanjing University and the balls were stainless steel. The mass ratio of ball to powder was 40:1 and that of large balls to small balls was 1:2. The rotation rate was 300 rpm and the milling time was 5 h. The particle size distribution of the ball milled fine powder was tested by a Nano ZetaSizer machine (Malvern, MS2000). The particle size ranged from 49.66 mm to 152.45 mm and the average size was about 87.28 mm. The hydrogenation thermodynamics of alloys are the foundation and basis of the hydrogenation experiments. From thermodynamics research, the process of hydrogenation in the alloy will be clarified and thermodynamics parameters such as temperature and hydrogen pressure will be obtained. According to the measurement for hydrogen storage materials, the hydrogen absorption and desorption behaviors of the Mg alloys could be examined using a custom-made Sievert type pressureecompositionetemperature (PeCeT) volumetric apparatus (repeatability errors < 3%) at various temperatures (T ¼ 623, 648, 673 and 698 K) which were controlled within 0.5 K. The PC-Isotherms of samples were obtained at their first hydrogen absorption/desorption cycle since the reaction observed in the absorption process of Mg-13Y alloy is not reversible. Maximum hydrogen pressure, minimum hydrogen pressure and sample weight are set as 4.0 MPa, 0.019 MPa and 0.5 g, respectively. The time interval for pressure variation (<0.0002 MPa) is set as 20 s, which means if the pressure variation remains under 0.0002 MPa for 20 s, this point with the pressure and the composition will be recorded. Since the materials in the present research were not used for hydrogen storage, the maximum hydrogen absorption was not taken into account. The phase transformations accompanying hydrogenation process were monitored by X-ray diffraction (XRD) using an apparatus (D/max 2550VL/PCX) equipped with a Cu Ka radiation source. An XRD holder for powder sample was designed to isolate the sample from air. The sample powder was sealed with scotch tape in the glove-box. Characterization of microstructure evolution in the bulk alloy was performed by an optical microscopy (OM, Zeiss, Axio Observer A1), a scanning electron microscopy (SEM, JEOL, JSM-7600F) equipped with an energy dispersive X-ray spectrometer (EDS) and a transmission electron microscopy (TEM, JEOL, JEM-2100). The rectangular tensile specimens with dimensions of 4 mm width, 1.5 mm thickness and 15 mm gauge length were tested at ambient temperature using a Zwick-20 kN material testing machine with a crosshead speed of 0.1 mm/min. An extensometer was used during mechanical testing, and the stress strain curves obtained were engineering stresseengineering strain. What is also worth mentioning is that all fine powders obtained from ball milling and used for PCT experiment and the subsequent XRD characterization are transferred under the oxygen-free circumstance. For bulk alloys, there is no oxidation and oxygen before and during the PCT experiments. However, they are inevitably exposed to air about one hour while preparing the OM, SEM and TEM specimens. And the XRD results of bulk alloys are also obtained in the air. First principle calculation was conducted by using the Materials studio 5.5 [16], which is based on the density
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
functional theory with a plane wave basis set. The total energies of Mg24Y5, Mg, YH2 and H2 were calculated by using ultrasoft pseudopotentials [17]. All calculations in the present study were done at T ¼ 0 K and taking the spin polarization into account [18].
is only around 8.4 kJ/mol H2 [22]. As solute element in the magnesium alloy, aluminum would scarcely change the hydrogenation plateau pressure [13]. The reaction enthalpy and entropy for Mg-13Y alloy were determined using the Van’t Hoff relationship: ln pH2 ¼
Results Hydrogenation of Mg-13Y powder alloy To investigate the hydrogenation process in the Mg-13Y alloy, the hydrogenation conditions, such as temperature and pressure, were examined. Because of the slow reaction rate, Mg-13Y powder samples were used to study the hydrogenation thermodynamics. The pressureecomposition isotherms of Mg-13Y alloy powders taken at different temperatures are shown in Fig. 1(a) and the corresponding absorption and desorption plateaus are shown in Table 1. As expected, the plateau pressure increased by increasing the temperature. However, the plateaus of Mg-13Y powders are higher than those of pure Mg at each temperature [19,20]. According to the research of Takamura et al. [13], the plateau regime of AZ31 alloy (particles less than 100 mm) at 623 K is 0.6 MPa which is similar to pure Mg [20] but much lower than 1.303 MPa of Mg13Y alloy. In MgeAl based alloys, aluminum is not inclined to form hydride [11e15,21] since the formation enthalpy of AlH3
8305
DH DS : RT R
From the PCT data of Mg-13Y powders, the Van’t Hoff plots (lnP vs. 1000/T) were drawn in Fig. 1(b). According to the fitting lines from the experimental data, the Van’t Hoff equation for Mg-13Y alloy powder was determined to be lnPab ¼ 5.10/ T þ 10.78 for hydrogen absorption and lnPde ¼ 8.84/T þ 16.03 for hydrogen desorption. The hydrogenation enthalpy and entropy were 42.40 2.90 kJ/mol H2 and 89.62 4.41 J/mol H2/K, while the dehydrogenation enthalpy and entropy were 73.51 3.42 kJ/mol H2 and 133.29 5.19 J/mol H2/K. The value of hydride formation enthalpy for Mg-13Y alloy powder is dramatically lower than that of the hydrogenation of pure Mg (74.7 kJ/mol H2) [23]. The differences of thermodynamics data indicate that the hydrogenation process of Mg-13Y alloy is quite different from that of pure Mg or MgeAl (-Zn) system alloy. The dehydrogenation enthalpy of Mg-13Y powder is quite close to that of MgH2, which indirectly demonstrates that only MgH2 decomposed during the dehydrogenation process. The discrepancies between the absorption and desorption plateau pressures decrease with the increasing of the temperature.
Fig. 1 e PCT curves of Mg-13Y alloy powders measured at different temperatures (a) and the corresponding Van’t Hoff plots (b). (c) is the PCT curve of Mg-13Y powder measured at 623 K. (d) is the XRD results for the ball milled Mg-13Y powder, the powder sample hydrogenated at 623 K under 0.37 MPa for 10 h, the powder sample hydrogenated at 623 K under 3.5 MPa for 10 h and the sample evacuated at 623 K for 2 h after the PCT process (T [ 623 K).
8306
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Table 1 e PCT data of Mg-13Y alloy powders at different temperatures. Temperature (K)
Small H2-absorption plateaus (MPa)
H2-absorption plateaus (MPa)
H2-desorption plateaus (MPa)
0.1109 0.2170 e e
1.3030 1.9364 2.4079 3.2132
0.6246 1.1344 1.7079 2.9668
623 648 673 698
The PCT (T ¼ 623 K and 648 K) curves show the small absorption plateaus because of the lower hydrogen absorption rates at the lower temperatures. The PCT (T ¼ 623 K) curve shown in Fig. 1(c) was chosen to investigate the hydrogenation process of Mg-13Y alloy powder. The other two samples were prepared at 623 K under 0.37 MPa and 3.5 MPa for 10 h, standing for the small absorption plateau and the absorption plateau, respectively. The phase transformations were characterized by XRD results which are shown in Fig. 1(d). The ball milled Mg-13Y alloy powder consists of Mg and Mg24Y5 phase. As is widely known, there would be some yttrium atoms dissolved in the a-Mg phase during solidification and ball milling processes. After the hydrogenation process performed at 623 K under 0.37 MPa for 10 h, the Mg24Y5 phase was hydrogenated into YH2 phase. However, when the ball milled powder was hydrogenated at 623 K under 3.5 MPa for 10 h, MgH2 phase and YH3 phase could be detected in addition to YH2 and Mg. According to the research of the hydriding behavior of Mg24Y5 compound [24], with increasing hydrogen pressure at 643 K, the disproportionation reaction yielded at 0.5 MPa metallic Mg þ YH2, then at 1 MPa Mg þ MgH2 þ YH2 þ YH3, and finally at 2 MPa a complete hydrogenation reaction occurred yielding an equilibrium mixture of only two phases MgH2 þ YH3. Therefore, the small plateau around 0.11 MPa shown in Fig. 1(c) is corresponding to the hydrogenation reaction:
Mg24Y5 þ 5H2 / 24Mg þ 5YH2
(1)
The enthalpy of reaction (1) is 194.6 kJ/mol H2 according to the first principle calculation, which demonstrates that Mg24Y5 is easier to be hydrogenated than Mg. The small plateau related to Mg24Y5 phase is not obvious at higher
temperatures (T ¼ 673 K, 698 K). This is probably because the quantity of Mg24Y5 phase in the Mg-13Y alloy decreases along with the increase of the temperature according to MgeY binary phase diagram. In addition, the hydrogen absorption plateau would rise along with the increase of the temperature. The increase of the hydrogen absorption rate at the higher temperature might also lead to the instability of the small plateau. At the main plateau around 1.3 MPa, a-Mg phase with the solute yttrium is hydrogenated and the YH2 is further hydrogenated into YH3. According to the XRD results of the Mg-13Y powder treated at 623 K under 0.37 MPa for 10 h and the sample evacuated at 623 K for 2 h after PCT process shown in Fig. 1(d), they both consist of Mg phase and yttrium hydrides. The relative integrated intensities of yttrium hydrides to Mg are 0.50 and 0.56, respectively, which indicates that the content of yttrium hydrides increases after the main hydrogenation plateau. The hydrogenation processes could be described as follow:
Mg(Y) þ H2 / Mg(Y)H2
(2)
2YH2 þ H2 / 2YH3
(3)
As shown in Fig. 1(d), after dehydrogenation process and evacuation at 623 K for 2 h, YH2 precipitated from Mg(Y)H2 firstly and the MgH2 transformed into Mg. According to the dehydrogenation enthalpy which is close to that of pure Mg, the main dehydrogenation process is MgH2 decomposed into Mg. It is noted that YH2 and YH3 phases did not decompose. The formation enthalpies for YH2 and YH3 are 210 [22] and 158 [22] kJ/mol H2. The high absolute values of enthalpies
Fig. 2 e XRD patterns of (a) the Mg-13Y bulk alloy and (b) the Mg-8Al bulk alloy before and after hydrogenation at 673 K under 4 MPa for 40 h.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
8307
Fig. 3 e Microstructure of the as-cast Mg-13Y bulk alloy: (a) optical micrograph; (b) SEM backscattered electron image.
show their high stability which means YH2 and YH3 are easy to form but difficult to decompose. Therefore, Mg-13Y alloy powder is not suitable for HDDR process. Yttrium in Mg24Y5 phase could be easily hydrogenated into YH2, which is not a reversible reaction [25e28] under the present treatment conditions.
Hydrogenation of Mg-13Y bulk alloy In order to hydrogenate the bulk alloy, relatively high temperature of 673 K and long reaction time of 40 h are required taking the poor hydrogenation kinetics into account. The phase transition during the hydrogenation process was observed by XRD. Fig. 2 shows the XRD patterns of Mg-13Y and Mg-8Al bulk alloys before and after hydrogenation at 673 K under 4 MPa for 40 h. As expected, in the Mg-13Y bulk alloy, some YH2 formed during the hydrogenation process. However, there is no MgH2 phase even though the hydrogen pressure is higher than 2.4 MPa, the H2 absorption plateau pressure of Mg-13Y powder at 673 K. A small amount of YH2 phase was found in the center of the hydrogenated Mg-13Y alloy, which indicates H2 has already diffused into the center of the bulk specimen after the hydrogenation process. Y2O3 phase was also found by XRD analysis in the hydrogenated alloy and this was probably caused by the oxidization of YH2 during the XRD detection in the air [25]. Compared to Mg-13Y alloy, MgH2 phase was detected from the Mg-8Al bulk alloy hydrogenated at the same condition. Mg solid-solution and Mg17Al12 phase
changed into MgH2 and Al3.16Mg1.84 phases with some remnant Mg. Aluminum hydride is more difficult to form than magnesium hydride according to the formation enthalpies. Therefore, magnesium including a-Mg phase in the Mg-8Al bulk alloy is firstly and easily hydrogenated into MgH2 [15]. The hydrogenation of magnesium in the bulk Mg alloy would lead to pulverization and brittleness of the matrix which is harmful to mechanical properties. Mg24Y5 phase in Mg-13Y bulk alloy is easier to be hydrogenated than Mg matrix, and thus the density of the bulk alloy would not be destroyed. Fig. 3 shows microstructure of as-cast Mg-13Y bulk alloy. The as-cast Mg-13Y alloy consists of a-Mg phase, fishboneshaped eutectic structure with Mg and Mg24Y5 phase, divorced eutectic Mg24Y5 phase and the cuboid phase. As reported in Refs. [29,30], the cuboid phase is believed to be dihydride phase of rare earth with face-centered cubic (fcc) structure. The cuboid phase in the present work is YH2 which forms during melting or casting. And this phase is not detected by XRD analysis because of its low volume fraction. The microstructure of Mg-13Y alloy hydrogenated at 673 K under 4 MPa for 40 h are shown in Fig. 4. Some coarse long lath-shaped precipitates appeared because of the long time heating. And the fishbone-shaped eutectic structure was destroyed by forming the fine cuboid phases. For a more detailed observation and phase identification in the hydrogenated Mg-13Y alloy, TEM results are shown in Fig. 5. The fishbone-shaped Mg24Y5 (bcc, a ¼ 1.1278 nm in PDF card) structure around point A in Fig. 5(a) is identified by the
Fig. 4 e SEM images of Mg-13Y alloy hydrogenated at 673 K under 4 MPa for 40 h.
8308
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Fig. 5 e (a) TEM bright-field image of Mg-13Y bulk alloy hydrogenated at 673 K under 4 MPa for 40 h; (b) TEM bright-field image of cuboid phase; (c) and (d) are the SAED patterns of the region around the points A and B with the electron beam// [113] and [011], respectively.
[113] zone-axis electron diffraction pattern shown in Fig. 5(c), to be body-centered cubic (bcc) with a ¼ 1.1258 nm. The brightfield TEM image of YH2 (fcc, a ¼ 0.5207 nm in PDF card) phase and its selected area electron diffraction (SAED) patterns taken along with [011] are shown in Fig. 5(b) and (d), demonstrating that the cuboid phase has a face-centered cubic (fcc) structure with lattice parameter of 0.5210 nm. The cuboid phases in Fig. 5(a) are all near the large Mg24Y5 phase. The TEM results indicate that the cuboid phases in Mg-13Y alloy are YH2 and that the large Mg24Y5 phases are destroyed by forming the fine YH2 cuboid phases. Table 2 shows the tensile properties of the as-cast Mg-13Y alloy, hydrogenated alloy treated at 673 K and 4 MPa H2 for 40 h and H2-free treated alloy tested at room temperature. The Mg-13Y alloy we used is a binary alloy and has not been optimized, and the casting process is also not the optimal. In addition, because of the small size of the sample chamber of PCT equipment, the tensile specimens have to be machined into non-standard ones. Therefore, the absolute values of the tensile testing results are not comparable with those of MgeRE alloys in other groups’ work. However, it is clear that the whole mechanical performance of Mg-13Y alloy heat treated at 673 K for 40 h is similar with that of the original ascast alloy, while the strengths and elongation of hydrogenated alloy (w140 MPa, w160 MPa and w1.5%) are all higher than those of as-cast alloy (w120 MPa, w130 MPa, w0.6%). It indicates that hydrogenation heat treatment might be an efficient method in improving the mechanical properties of MgeRE alloys. The improvement is mainly caused by the
Table 2 e Tensile properties of the as-cast Mg-13Y alloy, hydrogenated alloy treated at 673 K and 4 MPa H2 for 40 h and H2-free treated alloy tested at room temperature. Situation
TYS (MPa)
UTS (MPa)
Elongation (%)
As-cast 673 K/40 h 673 K/40 h/4 MPa
122.4 2.8 116.5 2.2 140.0 0.1
130.9 5.2 130.5 2.4 162.7 5.4
0.60 0.01 0.69 0.01 1.51 0.17
formation of fine cuboid YH2 phases as well as the dissolution of large Mg24Y5 second phases.
Discussion Hydrogenation mechanism of Mg-13Y powder alloy The hydrogenation enthalpy and entropy of Mg-13Y powder are abnormally lower than those of other Mg based materials while the dehydrogenation thermodynamics data of Mg-13Y powder are more similar with those of pure Mg. This phenomenon is mainly attributed to the solid solution of yttrium in the Mg matrix. It is widely known that there would be some yttrium dissolved in the a-Mg matrix in the Mg-13Y alloy during solidification and ball milling processes. This powder of Mg with some dispersively solute yttrium atoms is different from the ball milled mixture of Mg powder and Y powder [31] which has no solid solution effect. The hydrogenation process in the main absorption plateau is that a-Mg particles with yttrium solutes, Mg(Y), are hydrogenated into Mg(Y)H2. This reaction is different from the hydrogenation of pure Mg particles or the powder mixture of pure Mg and pure Y. Hydrogenation thermodynamics of Mg-13Y powder is quite different from dehydrogenation thermodynamics since main hydrogenation reactions and dehydrogenation reactions are totally different in the Mg-13Y alloy. The dehydrogenation process is mainly attributed to the decomposition of MgH2. In addition, the large particle size of around 100 mm and the inevitable oxide layer on the surface as well as detects in the alloy particles would also cause the slight increase of plateau pressures. As widely known that the powder with larger particles has a fewer surface area to interact with hydrogen. Therefore, another reason for the higher hydrogenation plateau pressure is that the system needs a higher pressure to force the hydrogen to break into and diffuse into the particles.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Hydrogenation mechanism of Mg-13Y bulk alloy After the hydrogenation process at a high hydrogen pressure of 4 MPa and at a high temperature of 673 K for a long reaction time of 40 h, the Mg-13Y bulk alloy was hydrogenated slightly because of the much lower hydrogen diffusion rate, much smaller surface area and much poorer hydrogenation kinetics. Only a part of Mg24Y5 phase was hydrogenated into YH2 and aMg phase did not change. However, a few YH2 phase existed in the center of bulk sample, which indicates that some hydrogen diffused into the center of the alloy. Compared to Mg matrix with yttrium solute, Mg24Y5 phase holds a higher hydrogenation enthalpy of around 194.6 kJ/mol H2 according to the calculation, and it is easier to be hydrogenated than Mg matrix. Therefore, Mg24Y5 phase would be firstly hydrogenated and in some sense protect Mg matrix from being hydrogenated. In addition, the plateau pressures of Mg matrix increased compared to pure Mg according to the PCT curves and thus the Mg phase in the Mg-13Y alloy is more difficult to be hydrogenated than in pure Mg. In the Mg-8Al bulk alloy, Mg atoms from both Mg17Al12 and a-Mg are easier to be hydrogenated than Al [11e15,21], and hence a large quantity of MgH2 appeared. In the bulk MgeRE alloy, Mg matrix could not transform into MgH2 while the large Mg24Y5 phases will be destroyed by forming the fine YH2 cuboid phases during the hydrogenation under 4 MPa H2 and under 673 K for not more than 40 h. The large eutectic structure consists of large Mg24Y5 phase and Mg will disappear gradually with some fine YH2 cuboid phases instead during the hydrogenation heat treatment. This may increase the mechanical properties of the thin-wall MgeRE casting alloy. The hydrogenation of magnesium in the bulk MgeAl alloy would probably lead to pulverization of the Mg matrix, the major part of the alloy, which is harmful to mechanical properties.
Conclusions The Mg-13Y casting alloy consists of Mg24Y5 phase and a-Mg with the yttrium solute, and these two phases have different hydrogenation enthalpies and plateau pressures according to the calculation result and PCT curves. Because of the poor diffusion and hydrogenation kinetics of the bulk alloy, the appropriate treatment parameters could be selected according to the powder research. On the other hand, the hydrogenation process and phase transitions were also investigated through the hydrogenation of Mg-13Y powder. In the first step of hydrogenation process, Mg24Y5 phase was hydrogenated into Mg and YH2 phase. After that the a-Mg with some yttrium solute began to transfer into Mg(Y)H2 and YH2 was further hydrogenated into YH3. However, in the dehydrogenation process, YH2 precipitated from Mg(Y)H2 and only MgH2 decomposed into Mg and H2, while the yttrium hydrides did not change. The abnormal hydrogenation thermodynamics data might be associated with the solid-solution pattern of yttrium in the Mg matrix. The cuboid phase in the hydrogenated Mg-13Y bulk alloy is identified as YH2 phase. The large Mg24Y5 phase could be
8309
destroyed into fine YH2 phase, which is probably the main cause for the improvement of mechanical properties of Mg13Y casting alloy. The Mg-Al based bulk alloy is not appropriate for hydrogenation heat treatment because the Mg could be hydrogenated into MgH2, which will destroy the Mg matrix by making it into powder.
Acknowledgments The authors gratefully acknowledge the financial supports of the National Natural Science Foundation of China (No. 51171113) and the Ministry of Science and Technology, China (No. 2011BAE22B02 and No. 2011 DFA50907).
references
[1] Shi XY, Li DJ, Luo AA, Hu B, Li L, Zeng XQ, et al. Microstructure and mechanical properties of Mg-7Al-2Sn alloy processed by super vacuum die-casting. Metall Mater Trans A 2013;44:4788e99. [2] Levi G, Avraham S, Zilberov A, Bamberger M. Solidification, solution treatment and age hardening of a Mg-1.6 wt.%Ca3.2 wt.%Zn alloy. Acta Mater 2006;54:523e30. [3] Gulyaev A. Physical metallurgy. Moscow: Mir Pub; 1980. [4] Hosmani SS, Schacherl RE, Mittemeijer EJ. Nitriding behavior of Fe-4 wt%V and Fe-2 wt%V alloys. Acta Mater 2005;53:2069e79. [5] Li W, Li XY, Dong HS. Effect of tensile stress on the formation of S-phase during low temperature plasma carburizing of 316L foil. Acta Mater 2011;59:5765e74. [6] Lagowski B. Metallography of hydrided ZE63 magnesium casting alloy. A F S Trans 1976;84:151. [7] Zhao ZY. High performance ZM8 casting magnesium Alloy processed by hydrogenation heat treatment (in Chinese). Aeronaut Mater 1979;1:1e7. [8] Takeshita T, Nakayama R. 10th Int workshop on rare earth magnets and their applications. Kyoto; 1989. pp. 551e62. [9] Harris IR, McGuiness PJ. Hydrogen: its use in the processing of NdFeB-type magnets. J Less Common Met 1991;172e174:1273e84. [10] Aoki K, Itoh Y, Park J, Kawamura Y, Inoue A, Masumoto T. Preparation of nanocrystalline bulk TiAl by a combination of HDDR and hot pressing. Intermetallics 1996;4:S103e11. [11] Hu LX, Wu Y, Yuan Y, Wang H. Microstructure nanocrystallization of a Mg-3 wt.%Al-1 wt.%Zn alloy by mechanically assisted hydriding-dehydriding. Mater Lett 2008;62:2984e7. [12] Sun HF, Fang W, Fang WB. Producing nanocrystalline bulk Mg-3Al-Zn alloy by powder metallurgy assisted hydridingdehydriding. J Alloy Compd 2011;509:8171e5. [13] Takamura H, Miyashita T, Kamegawa A, Okada M. Grain size refinement in Mg-Al-based alloy by hydrogen treatment. J Alloy Compd 2003;356e357:804e8. [14] Kamegawa A, Funayama T, Takahashi J, Takamura H, Okada M. Grain refinements of Al-Mg alloy by hydrogen heat-treatment. Mater Trans 2005;46:2449e53. [15] Miyazawa T, Kobayashi Y, Kamegawa A, Takamura H, Okada M. Grain size refinements of Mg Alloys (AZ61, AZ91, ZK60) by HDDR treatment. Mater Trans 2004;45:384e7. [16] Clark SJ, Segall MD, Pickard CJ, Hasnip PJ, Probert MJ, Refson K, et al. First principles methods using CASTEP. Z Fuer Kristallogr 2005;220:567e70.
8310
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
[17] Vanderbilt D. Soft self-consistent pseudopotentials in a generalized eigenvalue formalism. Phys Rev B 1990;41:7892e5. [18] Cheng LF, Zou JX, Zeng XQ, Ding WJ. Structure and thermodynamic studies of the LaMgNi4 compound and its hydrides by density functional theory. Intermetallics 2013;38:30e5. [19] Varin RA, Czujko T, Wronski ZS. Nanomaterials for solid state hydrogen storage. New York: Springer-Verlag New York Inc; 2008. [20] Phetsinorath S, Zou JX, Zeng XQ, Sun HQ, Ding WJ. Preparation and hydrogen storage properties of ultrafine pure Mg and Mg-Ti particles. Trans Nonferrous Met Soc China 2012;22:1849e54. [21] Zhang QA, Wu HY. Hydriding behavior of Mg17Al12 compound. Meter Chem Phys 2005;94:69e72. [22] Wolverton C, Ozolin‚s V, Asta M. Hydrogen in aluminum: first-principles calculations of structure and thermodynamics. Phys Rev B 2004;69:144109. [23] Aguey-Zinsou KF, Ares-Ferna´ndez JR. Hydrogen in magnesium: new perspectives toward functional stores. Energy Environ Sci 2010;3:526e43. [24] Zlotea C, Lu J, Andersson Y. Formation of one-dimensional MgH2 nano-structures by hydrogen induced disproportionation. J Alloy Compd 2006;426:357e62.
[25] Zhang QA, Liu DD, Wang QQ, Fang F, Sun DL, Ouyang LZ, et al. Superior hydrogen storage kinetics of Mg12YNi alloy with a long-period stacking ordered phase. Scr Mater 2011;65:233e6. [26] Sahlberg M, Andersson Y. Hydrogen absorption in Mg-Y-Zn ternary compounds. J Alloy Compd 2007;446e447:134e7. [27] Luo FP, Wang H, Ouyang LZ, Zeng MQ, Liu JW, Zhu M. Enhanced reversible hydrogen storage properties of a Mg-InY ternary solid solution. Int J Hydrogen Energy 2013;38:10912e8. [28] Liu JW, Zou CC, Wang H, Ouyang LZ, Zhu M. Facilitating de/ hydrogenation by long-period stacking ordered structure in Mg based alloys. Int J Hydrogen Energy 2013;38:10438e45. [29] Peng QM, Huang YD, Meng J, Li YD, Kainer KU. Strain induced GdH2 precipitate in Mg-Gd based alloys. Intermetallics 2011;19:382e9. [30] Yang YL, Peng LM, Fu PH, Hu B, Ding WJ. Identification of NdH2 particles in solution-treated Mg-2.5%Nd (wt.%) alloy. J Alloy Compd 2009;485:245e8. [31] Bystrzycki J, Czujko T, Varin RA. Processing by controlled mechanical milling of nanocomposite powders Mg þ X (X ¼ Co, Cr, Mo, V, Y, Zr) and their hydrogenation properties. J Alloy Compd 2005;404e406:507e10.
Available online at www.sciencedirect.com
ScienceDirect journal homepage: www.elsevier.com/locate/he
Study on hydrogenation behaviors of a Mg-13Y alloy Xiaoying Shi, Jianxin Zou, Chuan Liu, Lifang Cheng, Dejiang Li, Xiaoqin Zeng*, Wenjiang Ding The State Key Laboratory of Metal Matrix Composites, National Engineering Research Center of Light Alloy Net Forming, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, PR China
article info
abstract
Article history:
The Mg-13Y bulk alloy was prepared by conventional casting process and the Mg-13Y
Received 5 December 2013
powder was processed by ball milling using the casting alloy under the protection of
Received in revised form
argon. The hydrogenation thermodynamics, hydrogenation process and phase transitions
7 March 2014
were carefully investigated in the Mg-13Y powder alloy. It is shown that the Mg-13Y casting
Accepted 17 March 2014
alloy consists of Mg24Y5 phase and a-Mg containing yttrium which have different hydro-
Available online 22 April 2014
genation enthalpies, 195 kJ/mol H2 (by calculation) and 42 kJ/mol H2 (by Pressure eCompositioneTemperature experiment), respectively. The structure evolution and phase
Keywords:
transition in the Mg-13Y bulk alloy treated at 673 K and at 4 MPa for 40 h were observed by
Hydrogenation behavior
an optical microscopy (OM), a scanning electron microscopy (SEM), a transmission electron
Mg-13Y bulk alloy
microscopy (TEM) and X-ray diffraction (XRD). The large Mg24Y5 phase in the bulk Mg-13Y
Hydrogenation enthalpy
alloy could be destroyed into fine cuboid-shaped YH2 phases during the hydrogenation
YH2 cuboid phase
process, which is probably responsible for the improvement of mechanical properties of Mg-13Y alloy. Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Magnesium alloys, the lightest structural alloys developed so far, have great potential for lightweight applications, ranging from the portable electronic devices to automobile parts. However, the low strength limited the range of its application seriously. Therefore, lots of heat treatments, such as solution treatment [1] and ageing treatment [2], have been developed to improve its strength. Hydrogenation heat treatment applied to magnesium alloys is one of the thermo-chemical
treatments [3] which are not widely used. This is similar to nitriding [4] or carburizing [5] heat treatments implemented in steel and could be applied to thin-wall castings because hydrogen diffusion speed in bulk alloys is low [3]. Hydrogenation heat treatment was well applied to strengthen the MgeRE (Rare Earth)eZn casting alloys since the insoluble MgeREeZn phases at grain boundaries could be transformed into rare earth hydride and Zn atoms could dissolve into Mg matrix to enhance the strength of matrix. The newly formed rare earth hydrides exhibiting tiny particle-like morphology were thermodynamic stable and hard [6,7]. Large
* Corresponding author. Tel.: þ86 21 54742301; fax: þ86 21 34203730. E-mail addresses: [email protected] (X. Shi), [email protected] (J. Zou), [email protected] (C. Liu), chenglifang@china. com.cn (L. Cheng), [email protected] (D. Li), [email protected] (X. Zeng), [email protected] (W. Ding). http://dx.doi.org/10.1016/j.ijhydene.2014.03.115 0360-3199/Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
8304
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
improvements of ultimate tensile strength (UTS) (from 160 MPa to 292 MPa) and elongation (from 2.0% to 12.1%) compared to the as-cast alloy were observed in the ZM8 (Mg6 wt%Zn-2.5 wt%RE-0.5 wt%Zr) alloy treated at 753 K (480 C) and 1 atm pressure of hydrogen for 24 h. In addition, the method of hydrogen heat treatment followed with ageing treatment was more effective at strengthening this alloy. The UTS of aged alloy reached 316 MPa and the tensile yield strength (TYS) of aged alloy increased to 223 MPa from 129 MPa in the as-cast alloy [7]. Another important issue about hydrogenation which can not be ignored is that the method of hydrogenationedisproportionationedehydrogenationerecombination (HDDR) [8e12] could reduce the grain size of MgeAl alloys. In previous work, Takamura et al. [13e15] investigated grain size refinement in MgeAl-based alloy powders by hydrogenation treatment and found that the grain size of AZ31 alloy powders was reduced from 50 to 300 mm to about 100 nm in the case of the heat treatment at 623 K (350 C) under a hydrogen pressure of 7 MPa for 24 h. In addition, AZ31 alloy plate was also conducted HDDR process at 723 K (450 C) and its grain size was reduced to less than 500 nm, while the layer thickness being treated was limited to the range of about 20 mm from the surface after 36 h at 7 MPa pressure of H2 [13]. Mg matrix was hydrogenated into MgH2 in the HD process and then the MgH2 was dehydrogenated into Mg in the DR process. Microstructure and strengthening mechanism of hydrogenated MgeRE alloys were not investigated carefully and completely. It is significant to understand the microstructure evolution, phase transformation and corresponding thermodynamics during the hydrogenation process in order to make good use of this method and design technological parameter. In the present work, Mg-13Y binary alloy was selected as an example to investigate the microstructure evolution and strengthening mechanism of MgeRE alloys heat-treated in the hydrogen. Since the diffusion speed of hydrogen in the bulk alloy is very low, the hydrogenation process was firstly investigated in the Mg-13Y powder alloy. According to the Refs. [7] and [13], the hydrogenation behaviors of MgeRE based bulk alloy and MgeAl based bulk alloy are totally different. Therefore, the hydrogenation behaviors of Mg-13Y bulk alloy and Mg-8Al bulk alloy were also directly compared in this paper.
Experimental procedures Mg-13Y (Mg-12.85 wt% Y) and Mg-8Al (Mg-7.68 wt% Al) alloys were melted in steel crucible under a cover gas mixture of SF6 and CO2 using pure Mg, Mg-25Y (wt%) master alloy and pure Al. The melt was poured at 1003 K (730 C) into the mould preheated at 473 K (200 C). Bulk specimens with dimensions of 30 mm 10 mm 2 mm were cut from ingots by wire electrical discharge machining (WEDM) and removed oxide scale from the surface by using SiC paper in the Ar-filled glove box. The coarse powder samples used for thermodynamics researches were prepared in air by punching machine but quickly transferred into the Ar-filled glove box in approximately one minute. Under argon protection, the fine powder
samples were produced from coarse powder using a QM-3SP2 planetary ball milling machine made by Nanjing University and the balls were stainless steel. The mass ratio of ball to powder was 40:1 and that of large balls to small balls was 1:2. The rotation rate was 300 rpm and the milling time was 5 h. The particle size distribution of the ball milled fine powder was tested by a Nano ZetaSizer machine (Malvern, MS2000). The particle size ranged from 49.66 mm to 152.45 mm and the average size was about 87.28 mm. The hydrogenation thermodynamics of alloys are the foundation and basis of the hydrogenation experiments. From thermodynamics research, the process of hydrogenation in the alloy will be clarified and thermodynamics parameters such as temperature and hydrogen pressure will be obtained. According to the measurement for hydrogen storage materials, the hydrogen absorption and desorption behaviors of the Mg alloys could be examined using a custom-made Sievert type pressureecompositionetemperature (PeCeT) volumetric apparatus (repeatability errors < 3%) at various temperatures (T ¼ 623, 648, 673 and 698 K) which were controlled within 0.5 K. The PC-Isotherms of samples were obtained at their first hydrogen absorption/desorption cycle since the reaction observed in the absorption process of Mg-13Y alloy is not reversible. Maximum hydrogen pressure, minimum hydrogen pressure and sample weight are set as 4.0 MPa, 0.019 MPa and 0.5 g, respectively. The time interval for pressure variation (<0.0002 MPa) is set as 20 s, which means if the pressure variation remains under 0.0002 MPa for 20 s, this point with the pressure and the composition will be recorded. Since the materials in the present research were not used for hydrogen storage, the maximum hydrogen absorption was not taken into account. The phase transformations accompanying hydrogenation process were monitored by X-ray diffraction (XRD) using an apparatus (D/max 2550VL/PCX) equipped with a Cu Ka radiation source. An XRD holder for powder sample was designed to isolate the sample from air. The sample powder was sealed with scotch tape in the glove-box. Characterization of microstructure evolution in the bulk alloy was performed by an optical microscopy (OM, Zeiss, Axio Observer A1), a scanning electron microscopy (SEM, JEOL, JSM-7600F) equipped with an energy dispersive X-ray spectrometer (EDS) and a transmission electron microscopy (TEM, JEOL, JEM-2100). The rectangular tensile specimens with dimensions of 4 mm width, 1.5 mm thickness and 15 mm gauge length were tested at ambient temperature using a Zwick-20 kN material testing machine with a crosshead speed of 0.1 mm/min. An extensometer was used during mechanical testing, and the stress strain curves obtained were engineering stresseengineering strain. What is also worth mentioning is that all fine powders obtained from ball milling and used for PCT experiment and the subsequent XRD characterization are transferred under the oxygen-free circumstance. For bulk alloys, there is no oxidation and oxygen before and during the PCT experiments. However, they are inevitably exposed to air about one hour while preparing the OM, SEM and TEM specimens. And the XRD results of bulk alloys are also obtained in the air. First principle calculation was conducted by using the Materials studio 5.5 [16], which is based on the density
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
functional theory with a plane wave basis set. The total energies of Mg24Y5, Mg, YH2 and H2 were calculated by using ultrasoft pseudopotentials [17]. All calculations in the present study were done at T ¼ 0 K and taking the spin polarization into account [18].
is only around 8.4 kJ/mol H2 [22]. As solute element in the magnesium alloy, aluminum would scarcely change the hydrogenation plateau pressure [13]. The reaction enthalpy and entropy for Mg-13Y alloy were determined using the Van’t Hoff relationship: ln pH2 ¼
Results Hydrogenation of Mg-13Y powder alloy To investigate the hydrogenation process in the Mg-13Y alloy, the hydrogenation conditions, such as temperature and pressure, were examined. Because of the slow reaction rate, Mg-13Y powder samples were used to study the hydrogenation thermodynamics. The pressureecomposition isotherms of Mg-13Y alloy powders taken at different temperatures are shown in Fig. 1(a) and the corresponding absorption and desorption plateaus are shown in Table 1. As expected, the plateau pressure increased by increasing the temperature. However, the plateaus of Mg-13Y powders are higher than those of pure Mg at each temperature [19,20]. According to the research of Takamura et al. [13], the plateau regime of AZ31 alloy (particles less than 100 mm) at 623 K is 0.6 MPa which is similar to pure Mg [20] but much lower than 1.303 MPa of Mg13Y alloy. In MgeAl based alloys, aluminum is not inclined to form hydride [11e15,21] since the formation enthalpy of AlH3
8305
DH DS : RT R
From the PCT data of Mg-13Y powders, the Van’t Hoff plots (lnP vs. 1000/T) were drawn in Fig. 1(b). According to the fitting lines from the experimental data, the Van’t Hoff equation for Mg-13Y alloy powder was determined to be lnPab ¼ 5.10/ T þ 10.78 for hydrogen absorption and lnPde ¼ 8.84/T þ 16.03 for hydrogen desorption. The hydrogenation enthalpy and entropy were 42.40 2.90 kJ/mol H2 and 89.62 4.41 J/mol H2/K, while the dehydrogenation enthalpy and entropy were 73.51 3.42 kJ/mol H2 and 133.29 5.19 J/mol H2/K. The value of hydride formation enthalpy for Mg-13Y alloy powder is dramatically lower than that of the hydrogenation of pure Mg (74.7 kJ/mol H2) [23]. The differences of thermodynamics data indicate that the hydrogenation process of Mg-13Y alloy is quite different from that of pure Mg or MgeAl (-Zn) system alloy. The dehydrogenation enthalpy of Mg-13Y powder is quite close to that of MgH2, which indirectly demonstrates that only MgH2 decomposed during the dehydrogenation process. The discrepancies between the absorption and desorption plateau pressures decrease with the increasing of the temperature.
Fig. 1 e PCT curves of Mg-13Y alloy powders measured at different temperatures (a) and the corresponding Van’t Hoff plots (b). (c) is the PCT curve of Mg-13Y powder measured at 623 K. (d) is the XRD results for the ball milled Mg-13Y powder, the powder sample hydrogenated at 623 K under 0.37 MPa for 10 h, the powder sample hydrogenated at 623 K under 3.5 MPa for 10 h and the sample evacuated at 623 K for 2 h after the PCT process (T [ 623 K).
8306
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Table 1 e PCT data of Mg-13Y alloy powders at different temperatures. Temperature (K)
Small H2-absorption plateaus (MPa)
H2-absorption plateaus (MPa)
H2-desorption plateaus (MPa)
0.1109 0.2170 e e
1.3030 1.9364 2.4079 3.2132
0.6246 1.1344 1.7079 2.9668
623 648 673 698
The PCT (T ¼ 623 K and 648 K) curves show the small absorption plateaus because of the lower hydrogen absorption rates at the lower temperatures. The PCT (T ¼ 623 K) curve shown in Fig. 1(c) was chosen to investigate the hydrogenation process of Mg-13Y alloy powder. The other two samples were prepared at 623 K under 0.37 MPa and 3.5 MPa for 10 h, standing for the small absorption plateau and the absorption plateau, respectively. The phase transformations were characterized by XRD results which are shown in Fig. 1(d). The ball milled Mg-13Y alloy powder consists of Mg and Mg24Y5 phase. As is widely known, there would be some yttrium atoms dissolved in the a-Mg phase during solidification and ball milling processes. After the hydrogenation process performed at 623 K under 0.37 MPa for 10 h, the Mg24Y5 phase was hydrogenated into YH2 phase. However, when the ball milled powder was hydrogenated at 623 K under 3.5 MPa for 10 h, MgH2 phase and YH3 phase could be detected in addition to YH2 and Mg. According to the research of the hydriding behavior of Mg24Y5 compound [24], with increasing hydrogen pressure at 643 K, the disproportionation reaction yielded at 0.5 MPa metallic Mg þ YH2, then at 1 MPa Mg þ MgH2 þ YH2 þ YH3, and finally at 2 MPa a complete hydrogenation reaction occurred yielding an equilibrium mixture of only two phases MgH2 þ YH3. Therefore, the small plateau around 0.11 MPa shown in Fig. 1(c) is corresponding to the hydrogenation reaction:
Mg24Y5 þ 5H2 / 24Mg þ 5YH2
(1)
The enthalpy of reaction (1) is 194.6 kJ/mol H2 according to the first principle calculation, which demonstrates that Mg24Y5 is easier to be hydrogenated than Mg. The small plateau related to Mg24Y5 phase is not obvious at higher
temperatures (T ¼ 673 K, 698 K). This is probably because the quantity of Mg24Y5 phase in the Mg-13Y alloy decreases along with the increase of the temperature according to MgeY binary phase diagram. In addition, the hydrogen absorption plateau would rise along with the increase of the temperature. The increase of the hydrogen absorption rate at the higher temperature might also lead to the instability of the small plateau. At the main plateau around 1.3 MPa, a-Mg phase with the solute yttrium is hydrogenated and the YH2 is further hydrogenated into YH3. According to the XRD results of the Mg-13Y powder treated at 623 K under 0.37 MPa for 10 h and the sample evacuated at 623 K for 2 h after PCT process shown in Fig. 1(d), they both consist of Mg phase and yttrium hydrides. The relative integrated intensities of yttrium hydrides to Mg are 0.50 and 0.56, respectively, which indicates that the content of yttrium hydrides increases after the main hydrogenation plateau. The hydrogenation processes could be described as follow:
Mg(Y) þ H2 / Mg(Y)H2
(2)
2YH2 þ H2 / 2YH3
(3)
As shown in Fig. 1(d), after dehydrogenation process and evacuation at 623 K for 2 h, YH2 precipitated from Mg(Y)H2 firstly and the MgH2 transformed into Mg. According to the dehydrogenation enthalpy which is close to that of pure Mg, the main dehydrogenation process is MgH2 decomposed into Mg. It is noted that YH2 and YH3 phases did not decompose. The formation enthalpies for YH2 and YH3 are 210 [22] and 158 [22] kJ/mol H2. The high absolute values of enthalpies
Fig. 2 e XRD patterns of (a) the Mg-13Y bulk alloy and (b) the Mg-8Al bulk alloy before and after hydrogenation at 673 K under 4 MPa for 40 h.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
8307
Fig. 3 e Microstructure of the as-cast Mg-13Y bulk alloy: (a) optical micrograph; (b) SEM backscattered electron image.
show their high stability which means YH2 and YH3 are easy to form but difficult to decompose. Therefore, Mg-13Y alloy powder is not suitable for HDDR process. Yttrium in Mg24Y5 phase could be easily hydrogenated into YH2, which is not a reversible reaction [25e28] under the present treatment conditions.
Hydrogenation of Mg-13Y bulk alloy In order to hydrogenate the bulk alloy, relatively high temperature of 673 K and long reaction time of 40 h are required taking the poor hydrogenation kinetics into account. The phase transition during the hydrogenation process was observed by XRD. Fig. 2 shows the XRD patterns of Mg-13Y and Mg-8Al bulk alloys before and after hydrogenation at 673 K under 4 MPa for 40 h. As expected, in the Mg-13Y bulk alloy, some YH2 formed during the hydrogenation process. However, there is no MgH2 phase even though the hydrogen pressure is higher than 2.4 MPa, the H2 absorption plateau pressure of Mg-13Y powder at 673 K. A small amount of YH2 phase was found in the center of the hydrogenated Mg-13Y alloy, which indicates H2 has already diffused into the center of the bulk specimen after the hydrogenation process. Y2O3 phase was also found by XRD analysis in the hydrogenated alloy and this was probably caused by the oxidization of YH2 during the XRD detection in the air [25]. Compared to Mg-13Y alloy, MgH2 phase was detected from the Mg-8Al bulk alloy hydrogenated at the same condition. Mg solid-solution and Mg17Al12 phase
changed into MgH2 and Al3.16Mg1.84 phases with some remnant Mg. Aluminum hydride is more difficult to form than magnesium hydride according to the formation enthalpies. Therefore, magnesium including a-Mg phase in the Mg-8Al bulk alloy is firstly and easily hydrogenated into MgH2 [15]. The hydrogenation of magnesium in the bulk Mg alloy would lead to pulverization and brittleness of the matrix which is harmful to mechanical properties. Mg24Y5 phase in Mg-13Y bulk alloy is easier to be hydrogenated than Mg matrix, and thus the density of the bulk alloy would not be destroyed. Fig. 3 shows microstructure of as-cast Mg-13Y bulk alloy. The as-cast Mg-13Y alloy consists of a-Mg phase, fishboneshaped eutectic structure with Mg and Mg24Y5 phase, divorced eutectic Mg24Y5 phase and the cuboid phase. As reported in Refs. [29,30], the cuboid phase is believed to be dihydride phase of rare earth with face-centered cubic (fcc) structure. The cuboid phase in the present work is YH2 which forms during melting or casting. And this phase is not detected by XRD analysis because of its low volume fraction. The microstructure of Mg-13Y alloy hydrogenated at 673 K under 4 MPa for 40 h are shown in Fig. 4. Some coarse long lath-shaped precipitates appeared because of the long time heating. And the fishbone-shaped eutectic structure was destroyed by forming the fine cuboid phases. For a more detailed observation and phase identification in the hydrogenated Mg-13Y alloy, TEM results are shown in Fig. 5. The fishbone-shaped Mg24Y5 (bcc, a ¼ 1.1278 nm in PDF card) structure around point A in Fig. 5(a) is identified by the
Fig. 4 e SEM images of Mg-13Y alloy hydrogenated at 673 K under 4 MPa for 40 h.
8308
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Fig. 5 e (a) TEM bright-field image of Mg-13Y bulk alloy hydrogenated at 673 K under 4 MPa for 40 h; (b) TEM bright-field image of cuboid phase; (c) and (d) are the SAED patterns of the region around the points A and B with the electron beam// [113] and [011], respectively.
[113] zone-axis electron diffraction pattern shown in Fig. 5(c), to be body-centered cubic (bcc) with a ¼ 1.1258 nm. The brightfield TEM image of YH2 (fcc, a ¼ 0.5207 nm in PDF card) phase and its selected area electron diffraction (SAED) patterns taken along with [011] are shown in Fig. 5(b) and (d), demonstrating that the cuboid phase has a face-centered cubic (fcc) structure with lattice parameter of 0.5210 nm. The cuboid phases in Fig. 5(a) are all near the large Mg24Y5 phase. The TEM results indicate that the cuboid phases in Mg-13Y alloy are YH2 and that the large Mg24Y5 phases are destroyed by forming the fine YH2 cuboid phases. Table 2 shows the tensile properties of the as-cast Mg-13Y alloy, hydrogenated alloy treated at 673 K and 4 MPa H2 for 40 h and H2-free treated alloy tested at room temperature. The Mg-13Y alloy we used is a binary alloy and has not been optimized, and the casting process is also not the optimal. In addition, because of the small size of the sample chamber of PCT equipment, the tensile specimens have to be machined into non-standard ones. Therefore, the absolute values of the tensile testing results are not comparable with those of MgeRE alloys in other groups’ work. However, it is clear that the whole mechanical performance of Mg-13Y alloy heat treated at 673 K for 40 h is similar with that of the original ascast alloy, while the strengths and elongation of hydrogenated alloy (w140 MPa, w160 MPa and w1.5%) are all higher than those of as-cast alloy (w120 MPa, w130 MPa, w0.6%). It indicates that hydrogenation heat treatment might be an efficient method in improving the mechanical properties of MgeRE alloys. The improvement is mainly caused by the
Table 2 e Tensile properties of the as-cast Mg-13Y alloy, hydrogenated alloy treated at 673 K and 4 MPa H2 for 40 h and H2-free treated alloy tested at room temperature. Situation
TYS (MPa)
UTS (MPa)
Elongation (%)
As-cast 673 K/40 h 673 K/40 h/4 MPa
122.4 2.8 116.5 2.2 140.0 0.1
130.9 5.2 130.5 2.4 162.7 5.4
0.60 0.01 0.69 0.01 1.51 0.17
formation of fine cuboid YH2 phases as well as the dissolution of large Mg24Y5 second phases.
Discussion Hydrogenation mechanism of Mg-13Y powder alloy The hydrogenation enthalpy and entropy of Mg-13Y powder are abnormally lower than those of other Mg based materials while the dehydrogenation thermodynamics data of Mg-13Y powder are more similar with those of pure Mg. This phenomenon is mainly attributed to the solid solution of yttrium in the Mg matrix. It is widely known that there would be some yttrium dissolved in the a-Mg matrix in the Mg-13Y alloy during solidification and ball milling processes. This powder of Mg with some dispersively solute yttrium atoms is different from the ball milled mixture of Mg powder and Y powder [31] which has no solid solution effect. The hydrogenation process in the main absorption plateau is that a-Mg particles with yttrium solutes, Mg(Y), are hydrogenated into Mg(Y)H2. This reaction is different from the hydrogenation of pure Mg particles or the powder mixture of pure Mg and pure Y. Hydrogenation thermodynamics of Mg-13Y powder is quite different from dehydrogenation thermodynamics since main hydrogenation reactions and dehydrogenation reactions are totally different in the Mg-13Y alloy. The dehydrogenation process is mainly attributed to the decomposition of MgH2. In addition, the large particle size of around 100 mm and the inevitable oxide layer on the surface as well as detects in the alloy particles would also cause the slight increase of plateau pressures. As widely known that the powder with larger particles has a fewer surface area to interact with hydrogen. Therefore, another reason for the higher hydrogenation plateau pressure is that the system needs a higher pressure to force the hydrogen to break into and diffuse into the particles.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
Hydrogenation mechanism of Mg-13Y bulk alloy After the hydrogenation process at a high hydrogen pressure of 4 MPa and at a high temperature of 673 K for a long reaction time of 40 h, the Mg-13Y bulk alloy was hydrogenated slightly because of the much lower hydrogen diffusion rate, much smaller surface area and much poorer hydrogenation kinetics. Only a part of Mg24Y5 phase was hydrogenated into YH2 and aMg phase did not change. However, a few YH2 phase existed in the center of bulk sample, which indicates that some hydrogen diffused into the center of the alloy. Compared to Mg matrix with yttrium solute, Mg24Y5 phase holds a higher hydrogenation enthalpy of around 194.6 kJ/mol H2 according to the calculation, and it is easier to be hydrogenated than Mg matrix. Therefore, Mg24Y5 phase would be firstly hydrogenated and in some sense protect Mg matrix from being hydrogenated. In addition, the plateau pressures of Mg matrix increased compared to pure Mg according to the PCT curves and thus the Mg phase in the Mg-13Y alloy is more difficult to be hydrogenated than in pure Mg. In the Mg-8Al bulk alloy, Mg atoms from both Mg17Al12 and a-Mg are easier to be hydrogenated than Al [11e15,21], and hence a large quantity of MgH2 appeared. In the bulk MgeRE alloy, Mg matrix could not transform into MgH2 while the large Mg24Y5 phases will be destroyed by forming the fine YH2 cuboid phases during the hydrogenation under 4 MPa H2 and under 673 K for not more than 40 h. The large eutectic structure consists of large Mg24Y5 phase and Mg will disappear gradually with some fine YH2 cuboid phases instead during the hydrogenation heat treatment. This may increase the mechanical properties of the thin-wall MgeRE casting alloy. The hydrogenation of magnesium in the bulk MgeAl alloy would probably lead to pulverization of the Mg matrix, the major part of the alloy, which is harmful to mechanical properties.
Conclusions The Mg-13Y casting alloy consists of Mg24Y5 phase and a-Mg with the yttrium solute, and these two phases have different hydrogenation enthalpies and plateau pressures according to the calculation result and PCT curves. Because of the poor diffusion and hydrogenation kinetics of the bulk alloy, the appropriate treatment parameters could be selected according to the powder research. On the other hand, the hydrogenation process and phase transitions were also investigated through the hydrogenation of Mg-13Y powder. In the first step of hydrogenation process, Mg24Y5 phase was hydrogenated into Mg and YH2 phase. After that the a-Mg with some yttrium solute began to transfer into Mg(Y)H2 and YH2 was further hydrogenated into YH3. However, in the dehydrogenation process, YH2 precipitated from Mg(Y)H2 and only MgH2 decomposed into Mg and H2, while the yttrium hydrides did not change. The abnormal hydrogenation thermodynamics data might be associated with the solid-solution pattern of yttrium in the Mg matrix. The cuboid phase in the hydrogenated Mg-13Y bulk alloy is identified as YH2 phase. The large Mg24Y5 phase could be
8309
destroyed into fine YH2 phase, which is probably the main cause for the improvement of mechanical properties of Mg13Y casting alloy. The Mg-Al based bulk alloy is not appropriate for hydrogenation heat treatment because the Mg could be hydrogenated into MgH2, which will destroy the Mg matrix by making it into powder.
Acknowledgments The authors gratefully acknowledge the financial supports of the National Natural Science Foundation of China (No. 51171113) and the Ministry of Science and Technology, China (No. 2011BAE22B02 and No. 2011 DFA50907).
references
[1] Shi XY, Li DJ, Luo AA, Hu B, Li L, Zeng XQ, et al. Microstructure and mechanical properties of Mg-7Al-2Sn alloy processed by super vacuum die-casting. Metall Mater Trans A 2013;44:4788e99. [2] Levi G, Avraham S, Zilberov A, Bamberger M. Solidification, solution treatment and age hardening of a Mg-1.6 wt.%Ca3.2 wt.%Zn alloy. Acta Mater 2006;54:523e30. [3] Gulyaev A. Physical metallurgy. Moscow: Mir Pub; 1980. [4] Hosmani SS, Schacherl RE, Mittemeijer EJ. Nitriding behavior of Fe-4 wt%V and Fe-2 wt%V alloys. Acta Mater 2005;53:2069e79. [5] Li W, Li XY, Dong HS. Effect of tensile stress on the formation of S-phase during low temperature plasma carburizing of 316L foil. Acta Mater 2011;59:5765e74. [6] Lagowski B. Metallography of hydrided ZE63 magnesium casting alloy. A F S Trans 1976;84:151. [7] Zhao ZY. High performance ZM8 casting magnesium Alloy processed by hydrogenation heat treatment (in Chinese). Aeronaut Mater 1979;1:1e7. [8] Takeshita T, Nakayama R. 10th Int workshop on rare earth magnets and their applications. Kyoto; 1989. pp. 551e62. [9] Harris IR, McGuiness PJ. Hydrogen: its use in the processing of NdFeB-type magnets. J Less Common Met 1991;172e174:1273e84. [10] Aoki K, Itoh Y, Park J, Kawamura Y, Inoue A, Masumoto T. Preparation of nanocrystalline bulk TiAl by a combination of HDDR and hot pressing. Intermetallics 1996;4:S103e11. [11] Hu LX, Wu Y, Yuan Y, Wang H. Microstructure nanocrystallization of a Mg-3 wt.%Al-1 wt.%Zn alloy by mechanically assisted hydriding-dehydriding. Mater Lett 2008;62:2984e7. [12] Sun HF, Fang W, Fang WB. Producing nanocrystalline bulk Mg-3Al-Zn alloy by powder metallurgy assisted hydridingdehydriding. J Alloy Compd 2011;509:8171e5. [13] Takamura H, Miyashita T, Kamegawa A, Okada M. Grain size refinement in Mg-Al-based alloy by hydrogen treatment. J Alloy Compd 2003;356e357:804e8. [14] Kamegawa A, Funayama T, Takahashi J, Takamura H, Okada M. Grain refinements of Al-Mg alloy by hydrogen heat-treatment. Mater Trans 2005;46:2449e53. [15] Miyazawa T, Kobayashi Y, Kamegawa A, Takamura H, Okada M. Grain size refinements of Mg Alloys (AZ61, AZ91, ZK60) by HDDR treatment. Mater Trans 2004;45:384e7. [16] Clark SJ, Segall MD, Pickard CJ, Hasnip PJ, Probert MJ, Refson K, et al. First principles methods using CASTEP. Z Fuer Kristallogr 2005;220:567e70.
8310
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 9 ( 2 0 1 4 ) 8 3 0 3 e8 3 1 0
[17] Vanderbilt D. Soft self-consistent pseudopotentials in a generalized eigenvalue formalism. Phys Rev B 1990;41:7892e5. [18] Cheng LF, Zou JX, Zeng XQ, Ding WJ. Structure and thermodynamic studies of the LaMgNi4 compound and its hydrides by density functional theory. Intermetallics 2013;38:30e5. [19] Varin RA, Czujko T, Wronski ZS. Nanomaterials for solid state hydrogen storage. New York: Springer-Verlag New York Inc; 2008. [20] Phetsinorath S, Zou JX, Zeng XQ, Sun HQ, Ding WJ. Preparation and hydrogen storage properties of ultrafine pure Mg and Mg-Ti particles. Trans Nonferrous Met Soc China 2012;22:1849e54. [21] Zhang QA, Wu HY. Hydriding behavior of Mg17Al12 compound. Meter Chem Phys 2005;94:69e72. [22] Wolverton C, Ozolin‚s V, Asta M. Hydrogen in aluminum: first-principles calculations of structure and thermodynamics. Phys Rev B 2004;69:144109. [23] Aguey-Zinsou KF, Ares-Ferna´ndez JR. Hydrogen in magnesium: new perspectives toward functional stores. Energy Environ Sci 2010;3:526e43. [24] Zlotea C, Lu J, Andersson Y. Formation of one-dimensional MgH2 nano-structures by hydrogen induced disproportionation. J Alloy Compd 2006;426:357e62.
[25] Zhang QA, Liu DD, Wang QQ, Fang F, Sun DL, Ouyang LZ, et al. Superior hydrogen storage kinetics of Mg12YNi alloy with a long-period stacking ordered phase. Scr Mater 2011;65:233e6. [26] Sahlberg M, Andersson Y. Hydrogen absorption in Mg-Y-Zn ternary compounds. J Alloy Compd 2007;446e447:134e7. [27] Luo FP, Wang H, Ouyang LZ, Zeng MQ, Liu JW, Zhu M. Enhanced reversible hydrogen storage properties of a Mg-InY ternary solid solution. Int J Hydrogen Energy 2013;38:10912e8. [28] Liu JW, Zou CC, Wang H, Ouyang LZ, Zhu M. Facilitating de/ hydrogenation by long-period stacking ordered structure in Mg based alloys. Int J Hydrogen Energy 2013;38:10438e45. [29] Peng QM, Huang YD, Meng J, Li YD, Kainer KU. Strain induced GdH2 precipitate in Mg-Gd based alloys. Intermetallics 2011;19:382e9. [30] Yang YL, Peng LM, Fu PH, Hu B, Ding WJ. Identification of NdH2 particles in solution-treated Mg-2.5%Nd (wt.%) alloy. J Alloy Compd 2009;485:245e8. [31] Bystrzycki J, Czujko T, Varin RA. Processing by controlled mechanical milling of nanocomposite powders Mg þ X (X ¼ Co, Cr, Mo, V, Y, Zr) and their hydrogenation properties. J Alloy Compd 2005;404e406:507e10.