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Hydrogen in magnesium alloys and magnesium interfaces: preparation, electronic properties and interdiffusion
Journal of the Less-Common
808
Metals, 172-174
(1991) 808-815
Hydrogen in magnesium alloys and magnesium interfaces: preparation, electronic properties and interdiffusion A. Fischer, ~hys~kalisches
H. Kiistler and L. Schlapbach ~nstitut, Universit&t Fribourg,
CH-1700 Fribourg
~~W~tzeT~and~
Abstract We have evaporated atomic monolayers of magnesium on polycrystalline PdH, and observed, by means of photoelectron spectroscopy, interdi~usion of hydrogen and the formation of an MgH, overlayer at low temperatures. Above 170 “C hydrogen desorbs rapidly. After hydrogen desorption, interdiffusion occurs at the Mg-Pd interface. We also prepared magnesium-rich Mg-Li alloys and studied the absorption of hydrogen, in order to investigate possible metallization. Reactivity around room temperature is very slow; at 300-400 “C the Mg-Li alloy decomposes and MgH, is formed.
1. Introduction It is well known that some metal hydrides have interesting properties for energy technology and hydrogen storage. The low temperature hydrides of the LaNi,-, FeTi- and ZrMn,-type families reach desorption pressures of about 1 bar in the very useful temperature range from 10 to 50 “C; however, their hydrogen content does not exceed 2 wt.%. Nowadays substantial efforts are being made in research to find suitable metal hydrogen systems based on light metals such as magnesium. Magnesium forms a quite stable, electrically isolating, hydride MgHz which decomposes at 290 *C at 1 bar. The metal hydrogen bonding is mostly ionic [I], in agreement with the optical transparency [2]. The reaction Mg + H, + MgH, is complex and H, at pressures above some mbar, slow. This may have several different causes. Electronic surface properties of magnesium do not favour H, dissociation [3]. Sticking coefficients are estimated to be of the order of lo-” [4]. Another reason is the low diffusion rate of hydrogen in Mg and MgHz. Krozer and Kesemo proposed a passivating MgHz surface layer. The hydriding process is virtually stopped by the low diffusion rate in MgH, [5]. While an excellent attempt of Bogdanovich [6] solved the problem of low kinetics by precipitation of very fine active MgH, powder from a metal-organic solution, the inconvenience of the high thermodynamic stability of MgHz remains. We re-examined magnesium hydride with the following goals: to study the electronic structure of MgH,; to study effects of alloying of magnesium and the possibility of a metallic magnesium-based hydride, of lower stability.
~22-50~8~911$3.~
@ Elsevier Sequoia/Printed
in The Netherlands
809
It has been suggested [l] that MgH,, doped with a monovalent element could be a superconductor. We tried to hydride a series of magnesium samples doped with some percentage of lithium. The hydriding behaviour and some X-ray photoelectron spectroscopy (XPS) data are presented here. As our earlier attempts to analyse the electronic properties of MgH, failed because of very severe contamination (mostly segregation of bulk oxygen to the surface of MgH,) [2], we decided to study, using XPS, the possibility of the formation of MgH, from a thin magnesium overlayer evaporated on ol-PdH,, which acts both as a support and as a source of atomic hydrogen. First results are reported here.
2. Experimental
details
The surface analytical experiments were performed with a VG ESCALAB 5 photoelectron spectrometer at a base pressure below lo-i0 mbar. We used Mg Ka (1253.6 eV), Si Kcr (1740 eV) and He I (21.2 eV) radiation; the Au 4f,,, line was at 84.0 eV. XPS spectra have been numerically corrected for Kcr,, 4 satellites. The palladium samples were prepared by r.f. welding a 0.25 mm palladium foil on a nickel sample holder. They were heated to 400 “C and cleaned (sputtered) with Ar+ ions at 220.5 kV in repeated cycles. In XPS analysis no nickel was observed; the carbon contamination of the surface was below 1%. The oxygen contamination was estimated from the Pd 3~,,~,~,~ and 0 1s peaks to be below 3% and 7% for magnesium on palladium and PdH, respectively. To prepare samples of magnesium on palladium we cleaned the palladium substrate with 500 V Ar+ ions, and then evaporated at room temperature or at - 170 “C an overlayer of 20-40 A of magnesium. To prepare the samples of magnesium on PdH, we cleaned the palladium substrate, transferred it to the high pressure cell inside the ultrahigh vacuum (UHV) spectrometer and hydrided it at 1 bar H, pressure, T = 25 “C, for some 10 min. Then we cooled it to - 170 “C, released the I$ pressure, sputtered the sample again, and evaporated approximately 20 A of magnesium onto the cooled sample. Mg-Li alloys containing 4.7, 8.9, 16.8 and 30.1 at.% Li were prepared by enclosing appropriate amounts of magnesium (99.99%, Rare Earth Products Ltd.) and lithium (99O/,, Merck) in tantalum tubes in an argon glove-box and r.f. heating the sealed tubes in vacuum. X-ray diffraction showed single-phase Mg-Li solid solutions with lattice parameters from a =3.2066& c =5.201 A, to a = 3.1947 A, c =5.1411 A, for the 4.7 at.% Li and the 16.8 at.% Li samples respectively. Attempts to prepare Mg-Li alloys in iron crucibles revealed ferromagnetic alloys owing to iron dissolved as an impurity.
Mg on PdH,
Mg on Pd
bulk Mg 1309 -
1307
1305
1303
1301
BINDINGENERGY(eV)
Fig. 1. Mg 1s spectra of bulk magnesium (20 “C) and PdH, ( -170 “C) substrates.
and of magnesium
films evaporated
on palladium
Hydrogen absorption at 40 bar and ZOO-400 “C in a high pressure microbalance was extremely slow. For improved hydrogenation, samples were cut into turnings and further milled. For the XPS measurements Mg-Li samples were filed with a diamond file and scratched with a steel blade in UHV. The oxygen and carbon surface contamination was relatively high (25% and 7% respectively), especially after heating the sample (36% and 5% respectively).
3, Results
and discussion
3.1. ~a~n~siurn films on PdH, and palladium The X-ray photoelectron spectra of magnesium on palladium and magnesium on PdH, are shown in Fig. 1 (Mg 1s) and Fig. 2 (Pd 3d). For calibration we determined the Mg 1s peak of pure magnesium at 1303.5 eV. The Mg 1s peak of magnesium on palladium at room temperature exhibits no significant shift in binding energy relative to the bulk magnesium sample. In free electron metals the photoelectron process results not only in the emission of the photoelectron but also in the excitation of collective modes of the electrons, so-called plasmons, which appear as satellites in the spectra
811
,,-
:
.‘: _.._.I_
,...‘,..“..‘,,l’-.
..
;
‘-.>.,_:v
:
i, “_,_,
Mg on PdH,
Mg on Pd
bulk Pd _..’
.:
,_..,’ . -_,__--’ . ’
\_,>. ....-
‘.a.
345
343 -
341
339
337
335
PdH,
333
BINDING ENERGY (eV)
Fig. 2. Pd 3d spectra of palladium covered by 20 A magnesium films.
(20 C) and PdH,
( - 170 C) substrates
and of substrates
due to the related energy loss of the photoelectron. Six plasmon losses with an energy of 11.0 eV are seen in the spectra (not shown). All these plasmon loss peaks disappear completely in the magnesium evaporated on PdH,, indicating the disappearance of free-electron-like behaviour. In contrast to magnesium on palladium, the Mg 1s peak of magnesium on PdH, is doubled, with a main peak at 1305.4 and a weaker peak at 1303.5 eV (the magnesium metal peak). Apparently atomic hydrogen from the PdH, substrate diffused across the interface into the magnesium. There are two ways of interpreting the splitting. (1) First, this is a chemical shift, caused by hydrogen. Most of the magnesium is transformed into MgH,; some, probably at the interface, stays as magnesium. (2) Second, the MgH, layer is isolating and thus positively charged by the photoelectron process, which lowers the kinetic energy of the outgoing electrons. This leads to a satellite, shifted to higher binding energy. The
812
lowest layer is in contact with the conducting PdH,, and thus does not contribute to the shifted peak. As oxidation of magnesium causes core level shifts of about 1 eV only and hydrogenation shifts usually are smaller than oxidation shifts, the second explanation seems to be more realistic. Further investigations are planned to decide which explanation is correct. The Pd 3d spectra of magnesium on palladium show contrary to the Mg 1s peak a shift of 1.2 eV towards higher binding energies. This points to an alloying of palladium and magnesium at the’ interface. The remaining shoulders at about 335.2 eV and about 340.6 eV originate from bulk palladium (335.1 eV and 340.3 eV for Pd 3d,,, and Pd 3d,,, respectively). Again we see plasmon losses, now induced by the passage of Pd 3d electrons through the magnesium layer. In the spectra of magnesium on PdH, the Pd 3d peaks are only slightly shifted by about 0.1-0.2 eV compared with the bulk palladium spectra. This shift could be due to the formation of PdH at the interface (the shift of the Pd 3d,,, peak of PdH is 0.2 eV compared with palladium [7]). The valence band spectra support this interpretation. They will be published and discussed in detail elsewhere. The above results which are almost identical for evaporation and interface analysis at -170 “C and at room temperature respectively show that at the Mg-Pd interface interdiffusion occurs and an Mg--Pd interface alloy is formed. The interface alloy remains covered by magnesium. Krozer and Kasemo [8] did not observe any interdiffusion up to about 120 “C; however, they studied an interface much thicker than ours. On the contrary, magnesium does not form an interface Mg-Pd alloy at the Mg-PdH, interface, but an MgH, film, Hydrogen diffuses already at -170 “C from the PdH, substrate into the magnesium overlayer.
3.2. LIesorption of hydrogen from MgHz overlayers at higher temperatures In order to study the hydrogen desorption from the MgH, overlayer and the interdiffusion, we evaporated magnesium on a cooled PdH, sample ( -170 “C) and heated it stepwise to 230 “C over several hours. The spectra are shown in Fig. 3. At low temperatures up to about 50 “C, the spectra look exactly as described above. At 170 “C we observed a distinct increase in the H, pressure in our UHV system (up to lop7 mbar). At this temperature the intensity of the Mg Is peak at 1305.4 eV decreased and the peak at 1303.5 eV increased. After 25 min at 170 “C the 1303.5 eV peak dominated the spectra; only a weak shoulder remained at 1305.4 eV. The Pd 3d peaks shifted at 170 “C! from 335.2 eV and 340.4 eV to 336.3 eV and 341.7 eV respectively. The area under the curve of the Pd 3d peaks increased compared with the Mg 1s peak, according to an increase in the Pd:Mg ratio from 16% to 21%, within the probing depth of about 20 A. In the whole experiment no plasmon losses were observed. Our interpretation is the following. On the PdH, sample an MgH, layer was formed at low temperature. By heating, hydrogen desorption began. At
813
‘;
A_! -. c;_
I
1309
1307
1305
1303
1301
Fig. 3. Mg Is and Pd 3d spectra of magnesium function of increasing temperature.
170 “C the MgH, decomposed began and an inter-metallic was formed.
345
343
341
339
337
I
335
I
333
BINDINGENERGY(eV-) evaporated
at -170
C on PW,
substrates
as a
and interdiffusion of magnesium into palladium compound between magnesium and palladium
3.3. Mg-Li alloys We analysed the surface composition of an alloy of magnesium with 14.6 at.% Li by XPS (Fig. 4). After cleaning, the sample exhibited an Mg:Li ratio at the surface close to the nominal composition. Four plasmon losses were seen in the Mg 2s and 2p core level spectra. After heating the sample for 30 min to 150 “C, the Li 1s peak increased strongly, corresponding to a lithium enrichment up to 47% (12% Mg, 36% 0, 5% C). A comparison of the binding energy with the literature value [9] shows that most of the lithium atoms are present in an ionic compound, probably a lithium oxide or hydroxide. The magnesium atoms seem to form a metallic environment. We observe four plasmon losses for the magnesium peaks, but none for the Li 1s peak.
814
66
62 t
58
54
50
46
BINDING ENERGY (ev)
Fig. 4. Mg 2p and Li 1s spectra of magnesium after heating to 150 ‘C in UHV.
with 14.6% Li at room temperature
before and
This suggests that the lithium has segregated to the surface and the sample is covered with about 30 A of LiO, or Li+OH-. We cut the samples into fine turnings and exposed them to 40 bar H, pressure at temperatures from 200 to 400 “C. The samples sometimes did not exhibit any reaction for days, and then they suddenly began to absorb hydrogen. The more lithium they contained, the slower the hydriding was. The total amount of absorbed hydrogen decreased with increasing lithium content. For Mg, _,Li, H, we reached the following values: x = 4.7%, y = 1.8; x = 8.9%, y = 1.2; 3c > 14%, y < 1. For the sample Mg,,Li,, no hydrogen uptake was observed. X-ray diffraction analysis of the x = 4.7% sample revealed only MgH, lines. For increasing lithium content the diffraction lines of hexagonal magnesium became visible and stronger. In the sample with 30% Li, however, no MgH, lines were observed. Lattice parameters did not exhibit a significant change for the lithium-containing magnesium and MgH, phases. Thus we supporse that for elevated temperatures (more than 150 “C) the lithium segregates to the surface and builds a relatively thick, passivating layer of LiOH or LiOz. Only when this layer is cracked does the hydrogen uptake begin.
815
4. Conclusion UHV clean overlayers of MgH, can be formed by the evaporation of magnesium on PdH, substrates at temperatures low enough to prevent fast desorption of hydrogen from PdH,. The presence of hydrogen at the Pd-Mg interface prevents interdiffusion. Mg-Li solid solutions with up to 30 at.% Li do not react with H, gas at room temperature. At elevated temperature some reaction starts, but unfortunately decomposition with MgH, formation rather than metallization was observed.
Acknowledgments This work was financially supported by NEFF (Nationaler EnergieForschungs-Fonds). We also gratefully acknowledge the X-ray diffraction work done by K. Yvon, University of Geneva.
References 1 R. Yu and P. K. Lam, Phys. Reu. B, 37 (1988) 8730. 2 L. Schlapbach, J. Osterwalder and T. Riesterer, J. Less-Common Met., 103 (1984) 295. 3 J. K. Norskov, A. Houmoller, P. K. Johansson and B. I. Lundqvist, Phys. Reu. Lett., 46 (1981) 251. 4 L. Schlapbach, Surface properties and activation, In L. Schlapbach (ed.), Hydrogen in Zntermetallic Compounds ZZ, Topics in Applied Physics, Vol. 64, Springer, Berlin, in the press. 5 A. Krozer and B. Kasemo, J. Vuc. Sci. Z’echnol. A, 5 (1987) 1003. 6 B. Bogdanovic, Znt. J. Hydrogen Energy, 9 (1984) 937. 7 M. Gupta and L. Schlapbach, in L. Schlapbach (ed.), Hydrogen in Zntermetallic Compounds I, Topics in Applied Physics, Vol. 63, Springer, Berlin, 1988, Table 5.1. 8 A. Krozer, Ph.D. Thesis, Gothenburg, Sweden (1989). A. Krozer and B. Kasemo, J. Less-Common Met., (1990). to be published. 9 C. D. Wagner, W. M. Riggs, L. E. Davis, J. F. Moulder and G. E. Muilenberg (eds.), Handbook of X-Ray Photoelectron Spectroscopy, Perkin-Elmer, Eden Prairie, MN, 1979.
808
Metals, 172-174
(1991) 808-815
Hydrogen in magnesium alloys and magnesium interfaces: preparation, electronic properties and interdiffusion A. Fischer, ~hys~kalisches
H. Kiistler and L. Schlapbach ~nstitut, Universit&t Fribourg,
CH-1700 Fribourg
~~W~tzeT~and~
Abstract We have evaporated atomic monolayers of magnesium on polycrystalline PdH, and observed, by means of photoelectron spectroscopy, interdi~usion of hydrogen and the formation of an MgH, overlayer at low temperatures. Above 170 “C hydrogen desorbs rapidly. After hydrogen desorption, interdiffusion occurs at the Mg-Pd interface. We also prepared magnesium-rich Mg-Li alloys and studied the absorption of hydrogen, in order to investigate possible metallization. Reactivity around room temperature is very slow; at 300-400 “C the Mg-Li alloy decomposes and MgH, is formed.
1. Introduction It is well known that some metal hydrides have interesting properties for energy technology and hydrogen storage. The low temperature hydrides of the LaNi,-, FeTi- and ZrMn,-type families reach desorption pressures of about 1 bar in the very useful temperature range from 10 to 50 “C; however, their hydrogen content does not exceed 2 wt.%. Nowadays substantial efforts are being made in research to find suitable metal hydrogen systems based on light metals such as magnesium. Magnesium forms a quite stable, electrically isolating, hydride MgHz which decomposes at 290 *C at 1 bar. The metal hydrogen bonding is mostly ionic [I], in agreement with the optical transparency [2]. The reaction Mg + H, + MgH, is complex and H, at pressures above some mbar, slow. This may have several different causes. Electronic surface properties of magnesium do not favour H, dissociation [3]. Sticking coefficients are estimated to be of the order of lo-” [4]. Another reason is the low diffusion rate of hydrogen in Mg and MgHz. Krozer and Kesemo proposed a passivating MgHz surface layer. The hydriding process is virtually stopped by the low diffusion rate in MgH, [5]. While an excellent attempt of Bogdanovich [6] solved the problem of low kinetics by precipitation of very fine active MgH, powder from a metal-organic solution, the inconvenience of the high thermodynamic stability of MgHz remains. We re-examined magnesium hydride with the following goals: to study the electronic structure of MgH,; to study effects of alloying of magnesium and the possibility of a metallic magnesium-based hydride, of lower stability.
~22-50~8~911$3.~
@ Elsevier Sequoia/Printed
in The Netherlands
809
It has been suggested [l] that MgH,, doped with a monovalent element could be a superconductor. We tried to hydride a series of magnesium samples doped with some percentage of lithium. The hydriding behaviour and some X-ray photoelectron spectroscopy (XPS) data are presented here. As our earlier attempts to analyse the electronic properties of MgH, failed because of very severe contamination (mostly segregation of bulk oxygen to the surface of MgH,) [2], we decided to study, using XPS, the possibility of the formation of MgH, from a thin magnesium overlayer evaporated on ol-PdH,, which acts both as a support and as a source of atomic hydrogen. First results are reported here.
2. Experimental
details
The surface analytical experiments were performed with a VG ESCALAB 5 photoelectron spectrometer at a base pressure below lo-i0 mbar. We used Mg Ka (1253.6 eV), Si Kcr (1740 eV) and He I (21.2 eV) radiation; the Au 4f,,, line was at 84.0 eV. XPS spectra have been numerically corrected for Kcr,, 4 satellites. The palladium samples were prepared by r.f. welding a 0.25 mm palladium foil on a nickel sample holder. They were heated to 400 “C and cleaned (sputtered) with Ar+ ions at 220.5 kV in repeated cycles. In XPS analysis no nickel was observed; the carbon contamination of the surface was below 1%. The oxygen contamination was estimated from the Pd 3~,,~,~,~ and 0 1s peaks to be below 3% and 7% for magnesium on palladium and PdH, respectively. To prepare samples of magnesium on palladium we cleaned the palladium substrate with 500 V Ar+ ions, and then evaporated at room temperature or at - 170 “C an overlayer of 20-40 A of magnesium. To prepare the samples of magnesium on PdH, we cleaned the palladium substrate, transferred it to the high pressure cell inside the ultrahigh vacuum (UHV) spectrometer and hydrided it at 1 bar H, pressure, T = 25 “C, for some 10 min. Then we cooled it to - 170 “C, released the I$ pressure, sputtered the sample again, and evaporated approximately 20 A of magnesium onto the cooled sample. Mg-Li alloys containing 4.7, 8.9, 16.8 and 30.1 at.% Li were prepared by enclosing appropriate amounts of magnesium (99.99%, Rare Earth Products Ltd.) and lithium (99O/,, Merck) in tantalum tubes in an argon glove-box and r.f. heating the sealed tubes in vacuum. X-ray diffraction showed single-phase Mg-Li solid solutions with lattice parameters from a =3.2066& c =5.201 A, to a = 3.1947 A, c =5.1411 A, for the 4.7 at.% Li and the 16.8 at.% Li samples respectively. Attempts to prepare Mg-Li alloys in iron crucibles revealed ferromagnetic alloys owing to iron dissolved as an impurity.
Mg on PdH,
Mg on Pd
bulk Mg 1309 -
1307
1305
1303
1301
BINDINGENERGY(eV)
Fig. 1. Mg 1s spectra of bulk magnesium (20 “C) and PdH, ( -170 “C) substrates.
and of magnesium
films evaporated
on palladium
Hydrogen absorption at 40 bar and ZOO-400 “C in a high pressure microbalance was extremely slow. For improved hydrogenation, samples were cut into turnings and further milled. For the XPS measurements Mg-Li samples were filed with a diamond file and scratched with a steel blade in UHV. The oxygen and carbon surface contamination was relatively high (25% and 7% respectively), especially after heating the sample (36% and 5% respectively).
3, Results
and discussion
3.1. ~a~n~siurn films on PdH, and palladium The X-ray photoelectron spectra of magnesium on palladium and magnesium on PdH, are shown in Fig. 1 (Mg 1s) and Fig. 2 (Pd 3d). For calibration we determined the Mg 1s peak of pure magnesium at 1303.5 eV. The Mg 1s peak of magnesium on palladium at room temperature exhibits no significant shift in binding energy relative to the bulk magnesium sample. In free electron metals the photoelectron process results not only in the emission of the photoelectron but also in the excitation of collective modes of the electrons, so-called plasmons, which appear as satellites in the spectra
811
,,-
:
.‘: _.._.I_
,...‘,..“..‘,,l’-.
..
;
‘-.>.,_:v
:
i, “_,_,
Mg on PdH,
Mg on Pd
bulk Pd _..’
.:
,_..,’ . -_,__--’ . ’
\_,>. ....-
‘.a.
345
343 -
341
339
337
335
PdH,
333
BINDING ENERGY (eV)
Fig. 2. Pd 3d spectra of palladium covered by 20 A magnesium films.
(20 C) and PdH,
( - 170 C) substrates
and of substrates
due to the related energy loss of the photoelectron. Six plasmon losses with an energy of 11.0 eV are seen in the spectra (not shown). All these plasmon loss peaks disappear completely in the magnesium evaporated on PdH,, indicating the disappearance of free-electron-like behaviour. In contrast to magnesium on palladium, the Mg 1s peak of magnesium on PdH, is doubled, with a main peak at 1305.4 and a weaker peak at 1303.5 eV (the magnesium metal peak). Apparently atomic hydrogen from the PdH, substrate diffused across the interface into the magnesium. There are two ways of interpreting the splitting. (1) First, this is a chemical shift, caused by hydrogen. Most of the magnesium is transformed into MgH,; some, probably at the interface, stays as magnesium. (2) Second, the MgH, layer is isolating and thus positively charged by the photoelectron process, which lowers the kinetic energy of the outgoing electrons. This leads to a satellite, shifted to higher binding energy. The
812
lowest layer is in contact with the conducting PdH,, and thus does not contribute to the shifted peak. As oxidation of magnesium causes core level shifts of about 1 eV only and hydrogenation shifts usually are smaller than oxidation shifts, the second explanation seems to be more realistic. Further investigations are planned to decide which explanation is correct. The Pd 3d spectra of magnesium on palladium show contrary to the Mg 1s peak a shift of 1.2 eV towards higher binding energies. This points to an alloying of palladium and magnesium at the’ interface. The remaining shoulders at about 335.2 eV and about 340.6 eV originate from bulk palladium (335.1 eV and 340.3 eV for Pd 3d,,, and Pd 3d,,, respectively). Again we see plasmon losses, now induced by the passage of Pd 3d electrons through the magnesium layer. In the spectra of magnesium on PdH, the Pd 3d peaks are only slightly shifted by about 0.1-0.2 eV compared with the bulk palladium spectra. This shift could be due to the formation of PdH at the interface (the shift of the Pd 3d,,, peak of PdH is 0.2 eV compared with palladium [7]). The valence band spectra support this interpretation. They will be published and discussed in detail elsewhere. The above results which are almost identical for evaporation and interface analysis at -170 “C and at room temperature respectively show that at the Mg-Pd interface interdiffusion occurs and an Mg--Pd interface alloy is formed. The interface alloy remains covered by magnesium. Krozer and Kasemo [8] did not observe any interdiffusion up to about 120 “C; however, they studied an interface much thicker than ours. On the contrary, magnesium does not form an interface Mg-Pd alloy at the Mg-PdH, interface, but an MgH, film, Hydrogen diffuses already at -170 “C from the PdH, substrate into the magnesium overlayer.
3.2. LIesorption of hydrogen from MgHz overlayers at higher temperatures In order to study the hydrogen desorption from the MgH, overlayer and the interdiffusion, we evaporated magnesium on a cooled PdH, sample ( -170 “C) and heated it stepwise to 230 “C over several hours. The spectra are shown in Fig. 3. At low temperatures up to about 50 “C, the spectra look exactly as described above. At 170 “C we observed a distinct increase in the H, pressure in our UHV system (up to lop7 mbar). At this temperature the intensity of the Mg Is peak at 1305.4 eV decreased and the peak at 1303.5 eV increased. After 25 min at 170 “C the 1303.5 eV peak dominated the spectra; only a weak shoulder remained at 1305.4 eV. The Pd 3d peaks shifted at 170 “C! from 335.2 eV and 340.4 eV to 336.3 eV and 341.7 eV respectively. The area under the curve of the Pd 3d peaks increased compared with the Mg 1s peak, according to an increase in the Pd:Mg ratio from 16% to 21%, within the probing depth of about 20 A. In the whole experiment no plasmon losses were observed. Our interpretation is the following. On the PdH, sample an MgH, layer was formed at low temperature. By heating, hydrogen desorption began. At
813
‘;
A_! -. c;_
I
1309
1307
1305
1303
1301
Fig. 3. Mg Is and Pd 3d spectra of magnesium function of increasing temperature.
170 “C the MgH, decomposed began and an inter-metallic was formed.
345
343
341
339
337
I
335
I
333
BINDINGENERGY(eV-) evaporated
at -170
C on PW,
substrates
as a
and interdiffusion of magnesium into palladium compound between magnesium and palladium
3.3. Mg-Li alloys We analysed the surface composition of an alloy of magnesium with 14.6 at.% Li by XPS (Fig. 4). After cleaning, the sample exhibited an Mg:Li ratio at the surface close to the nominal composition. Four plasmon losses were seen in the Mg 2s and 2p core level spectra. After heating the sample for 30 min to 150 “C, the Li 1s peak increased strongly, corresponding to a lithium enrichment up to 47% (12% Mg, 36% 0, 5% C). A comparison of the binding energy with the literature value [9] shows that most of the lithium atoms are present in an ionic compound, probably a lithium oxide or hydroxide. The magnesium atoms seem to form a metallic environment. We observe four plasmon losses for the magnesium peaks, but none for the Li 1s peak.
814
66
62 t
58
54
50
46
BINDING ENERGY (ev)
Fig. 4. Mg 2p and Li 1s spectra of magnesium after heating to 150 ‘C in UHV.
with 14.6% Li at room temperature
before and
This suggests that the lithium has segregated to the surface and the sample is covered with about 30 A of LiO, or Li+OH-. We cut the samples into fine turnings and exposed them to 40 bar H, pressure at temperatures from 200 to 400 “C. The samples sometimes did not exhibit any reaction for days, and then they suddenly began to absorb hydrogen. The more lithium they contained, the slower the hydriding was. The total amount of absorbed hydrogen decreased with increasing lithium content. For Mg, _,Li, H, we reached the following values: x = 4.7%, y = 1.8; x = 8.9%, y = 1.2; 3c > 14%, y < 1. For the sample Mg,,Li,, no hydrogen uptake was observed. X-ray diffraction analysis of the x = 4.7% sample revealed only MgH, lines. For increasing lithium content the diffraction lines of hexagonal magnesium became visible and stronger. In the sample with 30% Li, however, no MgH, lines were observed. Lattice parameters did not exhibit a significant change for the lithium-containing magnesium and MgH, phases. Thus we supporse that for elevated temperatures (more than 150 “C) the lithium segregates to the surface and builds a relatively thick, passivating layer of LiOH or LiOz. Only when this layer is cracked does the hydrogen uptake begin.
815
4. Conclusion UHV clean overlayers of MgH, can be formed by the evaporation of magnesium on PdH, substrates at temperatures low enough to prevent fast desorption of hydrogen from PdH,. The presence of hydrogen at the Pd-Mg interface prevents interdiffusion. Mg-Li solid solutions with up to 30 at.% Li do not react with H, gas at room temperature. At elevated temperature some reaction starts, but unfortunately decomposition with MgH, formation rather than metallization was observed.
Acknowledgments This work was financially supported by NEFF (Nationaler EnergieForschungs-Fonds). We also gratefully acknowledge the X-ray diffraction work done by K. Yvon, University of Geneva.
References 1 R. Yu and P. K. Lam, Phys. Reu. B, 37 (1988) 8730. 2 L. Schlapbach, J. Osterwalder and T. Riesterer, J. Less-Common Met., 103 (1984) 295. 3 J. K. Norskov, A. Houmoller, P. K. Johansson and B. I. Lundqvist, Phys. Reu. Lett., 46 (1981) 251. 4 L. Schlapbach, Surface properties and activation, In L. Schlapbach (ed.), Hydrogen in Zntermetallic Compounds ZZ, Topics in Applied Physics, Vol. 64, Springer, Berlin, in the press. 5 A. Krozer and B. Kasemo, J. Vuc. Sci. Z’echnol. A, 5 (1987) 1003. 6 B. Bogdanovic, Znt. J. Hydrogen Energy, 9 (1984) 937. 7 M. Gupta and L. Schlapbach, in L. Schlapbach (ed.), Hydrogen in Zntermetallic Compounds I, Topics in Applied Physics, Vol. 63, Springer, Berlin, 1988, Table 5.1. 8 A. Krozer, Ph.D. Thesis, Gothenburg, Sweden (1989). A. Krozer and B. Kasemo, J. Less-Common Met., (1990). to be published. 9 C. D. Wagner, W. M. Riggs, L. E. Davis, J. F. Moulder and G. E. Muilenberg (eds.), Handbook of X-Ray Photoelectron Spectroscopy, Perkin-Elmer, Eden Prairie, MN, 1979.