The magnetocaloric effect, magnetic refrigeration and ductile intermetallic compounds

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Acta Materialia 57 (2009) 18–28 www.elsevier.com/locate/actamat

The magnetocaloric effect, magnetic refrigeration and ductile intermetallic compounds q K.A. Gschneidner Jr. Ames Laboratory, USDOE, and Department of Materials Science and Engineering, Iowa State University, 255 Spedding, Ames, IA 50011-3020, USA Received 10 July 2008; accepted 23 August 2008 Available online 26 September 2008

Abstract The magnetocaloric effect (MCE) is the ability of a magnetic material, particularly near its Curie temperature, to heat up when a magnetic field is applied (magnetization) and to cool when the field is removed (demagnetization). A number of new materials, especially Gd5(Si4xGex), have outstanding MCE properties near room temperature, significantly better than the well-known standard, Gd metal. The giant MCE effect observed in the Gd5(Si4xGex) alloys for x P 2 is due to a coupled magnetostructural transformation. The coupled transformation also accounts for several other interesting phenomena: giant magnetoresistance; colossal magnetostriction; spontaneous generation of voltage; unusual training, dynamical and thermal phenomena; acoustic emissions; and a novel glass-like kinetically retarded state. Extension of the work on the MCE has led to the near commercialization of magnetic refrigeration as a viable cooling technology, and to the discovery of ductile rare earth transition metal (RM) B2 CsCl-type intermetallic compounds. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Rare earth alloys; Crystal structure; Magnetic properties; Phase transformations; Mechanical properties

1. Introduction This paper highlights a few of the notable accomplishments that have resulted from collaborations with students, post-doctoral associates, colleagues and peers on a wide variety of topics over the last 25 years. Earlier accomplishments, before 1990, have been briefly summarized in my address upon winning the Frank H. Spedding Award at the 19th Rare Earth Research Conference [1]. Much of the research carried out in the late 1950s through the early 1990s, however, served as a strong foundation for more recent research activities, and helped to establish our expertise on the physical metallurgy and solid state physics of rare earth metals, alloys and compounds. Thus, because of our knowledge of rare earth magnetism Dr J. Barclay (Astronautics Corporation of America) asked us to develop a rare earth material which orders magnetically q

This Acta Materialia Gold Medal Lecture was presented on 10 March 2008 at the 137th Annual Meeting of TMS, New Orleans, USA. E-mail address: [email protected]

at 40 K and had as good magnetocaloric properties as a replacement of the expensive GdPd intermetallic compound which was being considered as a magnetic refrigerant for the liquefaction of hydrogen gas, which condenses at 20 K. Several new alloys were designed, the best of these was (Dy0.5Er0.5)Al2, which is not only less expensive (Al replacing Pd), but also had magnetocaloric properties which were 20% better than GdPd [2]. Shortly thereafter, an analysis by Barclay and myself indicated that near-room-temperature magnetic refrigeration could be competitive with conventional gas compression refrigeration for operating supermarket chillers and food processing plants with a 5-year payback in energy savings of 30% [2]. In addition, magnetic cooling is an environmentally friendly technology replacing greenhouse gases (CFCs and HFCs) or hazardous gases (NH3). This project was funded by the US Department of Energy’s Advanced Energy Projects group of the Office of Basic Energy Sciences. As a result of this research, two major breakthroughs were announced in 1997: (1) the successful demonstration that magnetic refrigeration was a viable

1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.08.048

K.A. Gschneidner Jr. / Acta Materialia 57 (2009) 18–28

cooling technology [2,3]; and (2) the discovery of the giant magnetocaloric effect (GMCE) in Gd5Si2Ge2 and related Gd5(Si4xGex) compounds [4,5]. These two discoveries set off a flurry of activities. Prior to 1997, ten to twenty papers per annum were published on the magnetocaloric effect (MCE), and 10 years later the number of papers published in 2007 was 250, more than a tenfold increase [2]. Also, before the successful demonstration that magnetic refrigeration was a viable cooling technology, only three near-room-temperature demonstration devices had been built but, by 2007, a total of 29 cooling units had been reported in the literature, with seven in 2007 alone [2]. The second breakthrough, however, changed the entire landscape, because it led to a whole new class of materials which exhibit extraordinary responses not only to weak triggers (magnetic fields) but also to strong triggers (temperature and pressure), which will be the subject of about two-thirds of this paper. 2. The Gd5(Si4xGex) extraordinary responsive magnetic materials The Gd5(Si4xGex) family of materials was discovered as a result of an applied project. In magnetic refrigeration, there was a need to find a magnetic refrigerant material which magnetically orders above room temperature so that the heat extracted from the object to be cooled can be rejected efficiently to the ambient. Gadolinium metal, which has a Curie temperature (TC) of 294 K, is the prototype magnetic refrigerant for near-room-temperature applications, but, unfortunately, it does not have a large enough MCE at 330 K in reasonably low magnetic fields (20 kOe). The compound Gd5Si4 was known to have a TC = 336 K, and that TC could be lowered by substituting Ge for Si [6]. Thus, we began a quantitative and more

Gd5(Si2.5Ge1.5) OI

50

extensive investigation of the influence of Ge on the TC of Gd5Si4. Ge substitutions in the Gd5(Si4xGex) pseudo binary system slowly lowered TC from 335 to 300 K as the Ge concentration approaches x = 2. At x = 2, the orthorhombic Gd5Si4-structure type (OI) [7] changed to a monoclinic distortion of the same, which is now known as the Gd5Si2Ge2-structure type (M) [7,8]. Heat capacity measurements revealed that the typical lambda-type second-order magnetic transition (paramagnetic (PM) to ferromagnetic (FM) upon cooling) for x < 2 (Fig. 1a) changed to a sharp delta function-like first-order transition peak when x = 2 (Fig. 1b). The transition in Gd5Si2Ge2 upon cooling is a coupled magnetostructural transition where the PM, M phase becomes FM, OI. From the heat capacity measured at zero and non-zero magnetic fields and from magnetization, one can calculate the MCE, and it is seen in Fig. 2 that, for Gd5Si2Ge2, it is about twice as large as that of Gd and Gd4Si2.5Ge1.5, and it was called the ‘‘giant MCE” (or GMCE) [4]. As this unusual behavior is associated with a crystal structure change, attention will now turn to the crystal structures of OI, M and the closely related Sm5Ge4-type structure (OII), which is observed in the Ge-rich phases of the Gd5(Si4xGex) alloys. 2.1. Crystal structure sequence in the Gd5(Si4xGex) pseudo binary system [7–9] The three closely related crystal structures [7–9] consist of self-assembled R5T4 nanolayers (slabs) which are either strongly interacting via pairs of Si, Ge atoms (which are ˚ ) or are weakly bonded together at a separation of 2.6 A interacting when these atoms are separated from each other ˚ . In the structure on the left-hand side at a distance of 3.5 A of Fig. 3 (OI), all the slabs are strongly and identically bonded and, below the magnetic ordering temperature, this

b H=0

20 kOe 50 kOe

40

30

20

10

0

Heat capacity, C (J mol-1 K-1)

Heat capacity, C (J mol-1 K-1)

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Gd5(Si2Ge2) M H=0

200

20 kOe 150

50 kOe

100

50

0 0

50

100 150 200 250 300 350

0

50

100 150 200 250 300 350

Temperature, T (K) Temperature, T (K) Fig. 1. The heat capacities of (a) Gd5(Si2.5Ge1.5), which exhibits a second-order magnetic transition, and (b) Gd5Si2Ge2, which exhibits a first-order magnetostructural transition, from 0 to 350 K in several magnetic fields.

K.A. Gschneidner Jr. / Acta Materialia 57 (2009) 18–28

Gd5(Si2.5Ge1.5) OI

10

Magnetocaloric effect, ΔTad (K)

a

8

ΔH=50 kOe

Gd ΔH=50 kOe

6

ΔH=20 kOe ΔH=20 kOe

4

2

b

Gd5(Si2Ge2) M

16

Magnetocaloric effect, ΔTad (K)

20

14

ΔH=50 kOe

12

Gd

10

ΔH=50 kOe

8

ΔH=20 kOe

6 4 2

0

ΔH=20 kOe

0 210

240

270

300

330

Temperature, T (K)

210

240

270

300

330

Temperature, T (K)

Fig. 2. The MCE in (a) Gd5(Si2.5Ge1.5) and (b) Gd5Si2Ge2, each compared with the MCE of Gd, as a function of temperature for magnetic field changes (DH) of 20 and 50 kOe.

Fig. 3. The three most important crystal structures, OI, M and OII, in the R5T4 family (where R = rare earth and T = Si, Ge, Sn, Pb, Ga, Sb) which give rise to the extraordinary behaviors to various external stimuli. The unit cells are shown as fine black lines in the three structures. The small spheres represent the Si, Ge atoms and the large spheres represent the Gd (or R) atoms. The strongly bonded Si, Ge pairs of atoms between the slabs (nanolayers) ˚ apart and are shown as connected with a heavy line, and for the weakly bonded Si, Ge pairs of atoms, which are 3.5 A ˚ apart, no line connects are 2.6 A ˚ in the weakly bonded and 5.4 A ˚ in the strongly bonded case. them. The next closest Si, Ge distances between the neighboring Si, Ge pairs are 4.1 A

phase is FM. This polymorph is known as the Gd5Si4-type structure. In the structure shown in the middle of Fig. 3 (M) pairs of slabs are strongly bonded together (the middle two nanolayers), but the adjacent slabs (top and bottom nanolayers) are weakly bonded to the middle pair of slabs. This polymorph, known as the Gd5Si2Ge2-type structure, is always PM in the Gd5T4 compounds, but not necessarily so for the R5T4 materials, when R is not Gd. The third structure of this sequence shown on the right-hand side of Fig. 3 is also orthorhombic (OII), but in this case all the slabs are weakly bonded by the Si, Ge pairs of atoms. This structure is either PM or antiferromagnetic (AFM), and has the Sm5Ge4-type structure. There is also a fourth structure, the Tm5Si2Sb2 type, which also has the same set of slabs,

but all the Si, Ge bond distances between the slabs are ˚ , i.e. there are no Si, Ge pairs of atoms equal at 4.1 A between the layers [9]. This fourth structure is only known when some antimony replaces Si or Ge and is not involved in the exotic behaviors observed in the R5T4 materials when T = Si and Ge. The transitions between OI and M and OI and OII take place by a shift of neighboring slabs along the a direction ˚ with respect to one another [8]. It is this shearing by 0.5 A which gives rise to the GMCE discussed above and also the colossal magnetostriction, giant magnetoresistance and spontaneous voltage generation behaviors reported for the Gd5(Si4xGex) alloys (see below). To date, no temperature, pressure or magnetic field induced transitions have

K.A. Gschneidner Jr. / Acta Materialia 57 (2009) 18–28

2.2. Other extraordinary magnetic responses Within a few years of the discovery of the GMCE in Gd5Si2Ge2, Morellon et al. reported that both Gd5Si1.8Ge2.2 [10] and Gd5(Si0.4Ge3.6) [11] exhibit an unusually large strain (1%) along the a axis when M transforms to OI [10] and OII transforms to OI [11], respectively, upon cooling. This linear striction was also confirmed by single crystal studies of Gd5Si2Ge2 [8], and it is of the order of 10,000 ppm, which is ten times larger than the magnetostriction in the well-known giant magnetostriction material Terfenol D [12]. Because of this large change in length during M ? OI transformation, the effect has been called ‘‘colossal” magnetostriction. It should be noted that the amount of magnetostriction increases gradually with increasing applied field in Terfenol D, while in Gd5Si1.8Ge2.2, once the magnetic field is strong enough to trigger the M to OI transition, the colossal magnetostriction occurs at once, and a further increase in the field has little or no effect on the magnetostriction. The discovery of the giant magnetoresistance in the Gd5(Si4xGex) alloys followed shortly thereafter, as reported independently by the Spanish [13] group and the USA group [14,15]. The giant magnetoresistance (25%) is about the same as that found in artificial multi-layered thin films (for which A. Fert and P. Gru¨nberg were awarded the 2007 Nobel Prize in physics [16]). The giant magnetoresistance in the Gd5(Si4xGex) alloys, however, can be either positive or negative, depending upon the Ge to Si ratio. In addition to these unusually strong responses to the magnetic field or temperature changes, which cause the M ? OI (field increase, temperature decrease) or OI ? M (field decrease, temperature increase) transformation, Levin et al. [17] observed a spontaneous voltage generation on passing through these transformations, when the resistance probes were attached to the sample, and no current was being applied to the Gd5Si1.95Ge2.05 sample. Similar behaviors were also observed in other Gd5(Si4xGex) alloys for x = 2.5 and 2 for the same transformation, and for x = 3.67 and 4.0, where the first-order transition is from OII to OI upon cooling or upon magnetization. Several years later, Zou et al. [18] found that this generated voltage was slightly anisotropic in a study of single crystals of Gd5Si2Ge2. 2.3. X-ray magnetic circular dichroism

˚ ) between the two slabs in the OI FM phase, which (2.7 A disappears when Ge atom pairs are no longer bonding (i.e. ˚ ) in the M (or OII) the Ge–Ge distance increases to 3.5 A PM phase [19]. The induced magnetic moment of 0.06lB on the 4p electrons of the Ge atoms was estimated from the density of states curves reported by Paudyal et al. [20], and is due to hybridization with the sd electrons of the Gd atoms [19,20]. This hybridization enables the RKKY (Ruderman–Kittel–Kasuya–Yosida) FM interactions between 4f Gd moments in the neighboring slabs. That is, the magnetic behavior of the R5T4 compounds is driven by the magnetism of the nanoslabs and by the Ge(Si) atom pairs connecting the nanoslabs. 2.4. Gd5Ge4 The Ge end member of the Gd5(Si4xGex) pseudo binary system is the crown jewel of the R5T4 family, exhibiting many exciting and unusual features which are related to the instability of its OII crystal structure as the temperature approaches 0 K. As one cools down from room temperature, the PM state becomes AFM at 128 K, and the compound retains this magnetic structure and also the OII crystal structure down to 2 K (lower temperature limit of the experimental apparati) in a zero magnetic field [9]. The Gd 4f magnetic moments in each slab are aligned ferromagnetically along the c axis, but the FM aligned moments in the neighboring slabs are aligned antiparallel, thus giving rise to an unusual long-range AFM structure [21]. The AFM state can be transformed to the OI(FM) state if the applied magnetic field is >10 kOe. When the Gd5Ge4 phase is demagnetized, it remains in the FM state (i.e. the transformation is irreversible) as long as the temperature is <10 K. If a zero-field cooled specimen is magnetized and demagnetized above 20 K the transformation

OI

OII

OII

FM

AFM

PM

40

KA state 20

SRFMC GP

0 0

Haskel et al. [19] carried out X-ray magnetic circular dichroism (XMCD) investigation of Gd5(Si2Ge2) and Gd5(SiGe3) by studying the change of the intensity at the Ge K edge and Gd L3 edge going through the magnetostructural change of these two compounds. They found a magnetic moment on the bonding Ge atom pairs

TN

H || b 60

Magnetic field, H (kOe)

been observed between M and OII, although one might expect such to occur.

21

50

100

150

20 200

250 50

Temperature, T (K) Fig. 4. The magnetic phase diagram for Gd5Ge4 when the magnetic field is applied parallel to the b axis at low temperatures (0–250 K) and low magnetic field (0–60 kOe). The acronyms KA, SRFMC and GP mean kinetically arrested, short-range ferromagnetic correlations and Griffithlike phases, respectively.

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K.A. Gschneidner Jr. / Acta Materialia 57 (2009) 18–28

becomes completely reversible behaving like a typical metamagnetic material. Between 10 and 20 K the transformation is partially reversible. As noted by Roy et al. [22,23] in addition to the OII(AFM) M OI(FM) transition phase boundary (see the H vs T magnetic phase diagram, Fig. 4), there is a freezing–unfreezing, glass-like transition which partially or completely kinetically arrests this transformation (see Fig. 4). Part of the sluggishness of the OII to O(1) (and reverse) transformation may be due to large shifts of the slabs in the a direction, which leads to 1.9% strain [24], approximately twice that observed between the OI and M structure in Gd5Si2Ge2 (see Section 2.1). Also as shown in Fig. 4, there are short-range dynamic FM correlations, i.e. FM clusters in the OII AFM phase above 70 K up to the Nee´l temperature of 128 K [24,25]. Above the AFM–PM phase boundary, Ouyang et al. [25] also observed FM clusters in the PM phase (i.e. a Griffith-like phase [26]). Magnetic fields >5 kOe will destroy these FM correlations in both the AFM matrix and the PM phase, and above 240 K the thermal energy is sufficient to prevent the Griffith-like FM clusters from forming. The formation of the Griffiths-like phase is thought to be due to the competition of the interslab and intraslab magnetic exchange interactions, and may be enhanced by the presence of the thin Gd5Ge3 platelets present in the sample [27]. 3. Other R5(Si4xGex) systems In view of the unusual extremum properties observed in the Gd5(Si4xGex) pseudo-binary system alloys many scientists, including ourselves, began to study other R5T4 systems where R = rare earths and T = Si, Ge, Sn, Pb, Ga, Sb. Because the Gd 4f charge density is spherical, crystal field effects are small and little or no anisotropic behavior is observed. But for the other lanthanides, the 4f electron densities are anisotropic, and thus crystalline electric field effects are normally quite pronounced, and so some unex-

a

pected and remarkable magnetic, electrical and thermal effects and behavior can be expected in these non-Gd R5T4 phases. Below are highlighted some of the novel and noteworthy discoveries that were unraveled over the last five years. For the other (other than Gd) R elements, in addition to the three nano-layered structures (OI, OII and M) the R5Si4 compounds containing the larger light lanthanides (La, Ce, Pr and Nd) crystallize in the tetragonal Zr5Si4 structure, but the corresponding R5Ge4 phases form the OII, Sm5Ge4-type structure [9,28]. The Zr5Si4 structure does not contain two-dimensional slabs, nor is it layered and, therefore, no unusual phenomena are expected unless, at the appropriate Si:Ge ratio, this phase transforms to OI, M or OII. The lattice parameters for the R5Si4 OI and the R5Ge4 OII compounds, and for the R5(Si4xGex) M phases where x ffi 2, are shown in Figs. 5a, b and 6, respectively. The lattice parameter vs atomic number plots of Fig. 5 exhibit the usual lanthanide contraction, and there is only one anomaly, at Yb. This indicates that all the lanthanide elements are trivalent in the R5Si4 and RGe4 compounds, except for Yb, which has a valence of 2.5. The Yb5(Si4xGex) phases are discussed in more detail in Section 3.4. It should be noted that Yb5Ge4 has the OI structure and not the OII structure that is observed for the other R5Ge4 compounds (Fig 5b), and this accounts for the negative deviation of the a lattice parameter from the straight line established by the other R5Ge4 phases. The lattice parameters and unit cell volume for the monoclinic phases which have Gd5Si2Ge2-type structure are shown in Fig. 6. As M phases form at various Si:Ge ratios because of size restrictions, the Si content for the light lanthanides progressively changes to larger Si values for the smaller heavy lanthanides; the Si and Ge concentration values are listed next to the data points. Again, the four plots exhibit the normal lanthanide contraction. Because Yb does not form the M phase, no anomaly is observed. It is noted that lattice parameters for the

b

15.5

15.5

Y 15.0

c 14.5

Zr5Si4-type

b

8.1

OI R5Si4

a

7.8 7.5

c a

7.2

58

60

62

64

66

Atomic number

68

70

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Lattice parameter (Å)

15.0

14.5

OII R5Ge4

b

8.1

Yb5Ge4 has OI

7.8

c a

7.5 7.2

58

60

62

64

66

68

70

Atomic number

Fig. 5. The lattice parameters as a function of atomic number (a) for the Zr5Si4 tetragonal-type and the Gd5Si5 OI-type structures and (b) for the Sm5Ge4 OII-type structure. The Yb5Ge4 phase has the OI-type structure (see text).

K.A. Gschneidner Jr. / Acta Materialia 57 (2009) 18–28

a

23

b

8.2

1050 15.2

7.0

3.5 : 0.5

900 ? DNE DNE

3.2 : 0.8

3.5 : 0.5 (Y) 2:2

2:2

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14.2

950

3:1

? DNE DNE

7.2

DNE

2:2

V (Å3)

14.4

1000

2:2

DNE

1.5 : 2.5

3.2 : 0.8

3.5 : 0.5

3.5 : 0.5 (Y)

14.8

3:1

a (Å)

2:2

7.4

b (Å) 15.0

14.6 2:2

2:2

7.6

2:2 2:2

DNE

1.5 : 2.5

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La Ce Pr Nd Pm Sm Eu Gd Tb Dy Ho Er Tm Yb Lu

Atomic Number

850

57

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60

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64

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68

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La Ce Pr Nd Pm Sm Eu Gd Tb Dy Ho Er Tm Yb Lu

Atomic Number

Fig. 6. The lattice parameters of the monoclinic phases in the R5(Si4xGex) systems. The a and c lattice parameters are given in (a), and the b parameter and unit cell volume V are plotted in (b). The numerical values listed next to the data points are the Si:Ge ratio for the given R. The Y5(Si3.5Ge0.5) values are plotted between Tb and Dy. The letters DNE for a given R means that the M phase ‘‘Does Not Exist” in the R5(Si4xGex) pseudo binary system.

Y5Si3.5Ge0.5 compound is included in Fig. 6, and it is plotted between Tb and Dy, which is the typical location for Ycontaining compounds in such plots. For R = Eu, no Eu5T4 phases have been prepared by the standard techniques. It is possible that the Eu5T4 compounds might be prepared by a high-pressure synthesis technique or by ball milling but, as far as I am aware no one has attempted to prepare a Eu 5:4 phase by these techniques. The Lu5(Si4xGex) compounds only form the OII phase for any value of x, presumably Lu is too small to form the OI structure. This is consistent with the trends in the other heavy lanthanides where the OII phase extends to smaller and smaller x values as the atomic number of the lanthanide metal increases, i.e. x = 2.8 for Gd decreases to x = 1.2 for Er [9]. For Y5(Si4xGex) pseudo binary system, the OII extends from x = 4 to x = 1.2, and the M phase exists from x = 0.6 to x = 0 [9]. The existence of the M phase in the pseudo-binary system suggests that the relationship between OI, M and OII is governed by the usual alloying theory criteria of size, valence and electronegativities, and is not just due to the magnetic contribution to phase stability. In its alloying behavior, Y is quite similar to that of Dy for most compounds and solid solutions alloys but, for the Dy5(Si4xGex) alloys, OII exists from x = 4 to x = 1.5, the M phase from 1.5 to 1.2 and OI from 1.2 to 0 [9]. Thus, it would appear that the magnetic contribution to phase stability favors the M phase relative to OII, and the OI structure relative to M. Below, some unusual and exotic behaviors in the R5T4 systems for R = Sm, Tb, Er and Yb are discussed. The systems for R = La, Ce, Pr, Nd, Dy, Ho and Tm are currently being investigated, and so these seven R5T4 systems will not be discussed. 3.1. Sm5(Si4xGex) system The Sm ion in the Sm5(Si4xGex) system is trivalent (see Section 3 and Figs. 5 and 6) and not mixed valent

(2 < v < 3) as in many Sm systems, nor is it divalent. This in itself is not too unusual, the interesting fact is that, in contrast to the behavior observed in most of the R5T4 systems, the magnetic fields up to 100 kOe essentially have no effect on the magnetic ordering temperatures (Fig. 7) [29]. In contrast, in Gd5(Si2Ge2) at 100 kOe, the Curie temperature is increased by 55 K from 270 to 325 K (Fig. 7d). The 100 kOe magnetic field, however, decreases the value of the heat capacity at the Curie temperature by 10% in Sm5Si4 (Fig. 7a), 0% in both Sm5Si2Ge2 (Fig. 7b) and Sm5Ge4 (Fig. 7c). As a result, there is little or no MCE in the Sm5(Si4xGex) alloys at magnetostructural transformations. Furthermore, the magnetic ordering temperatures for the Sm5(Si4xGex) compounds are much higher (by 150 K) than those of the corresponding light lanthanide compounds, as might be expected from the de Gennes factor [30]. This anomalous behavior is not yet understood. 3.2. Tb5(Si4xGex) system The Tb5(Si4xGex) system is second most studied of the R:Si(Ge) 5:4 compounds after the Gd 5:4 alloys. The crystal structures in the Tb:Si(Ge) 5:4 compounds are identical to those observed in the Gd5(Si4xGex) system, but the composition ranges (Si:Ge ratio) of the OII and M phases extend to smaller x values in the Tb alloys. The magnetic structures [31,32] in Gd5T4 and Tb5T4 alloys are similar in that the magnetic spins are aligned parallel (ferromagnetically) in Gd and nearly so in Tb in the slabs (ac plane), and the magnetic ordering behavior is dependent on whether or not the slabs couple ferromagnetically or antiferromagnetically in the b direction. The nearly FM arrangement of the Tb moments in the ac planes are slightly canted out of the ac plane. Also, there is a spin reorientation in the Tb materials at 50 K below the magnetic order–disorder transition, which is not observed in the corresponding Gd compounds. Another major difference is that, just above this magnetic ordering temperature,

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K.A. Gschneidner Jr. / Acta Materialia 57 (2009) 18–28 40

Sm5Si4 OI

Heat capacity, CP (J/ g-at K)

Heat capacity, CP (J/ g-at K)

40

30

20

10

H=0 Hμ0=0

Sm5Si2Ge2 M 30

20

10

H=0 Hμ0=0 H = 10kOe T Hμ0=100

H = 10kOe T Hμ0=100 0

0 0

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350

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350

Temperature, T (K)

Temperature, T (K)

240

Sm5Ge4 OII 30

20

10

H H=0 =0 H=100 = 100 kOe H kOe 0

Heat capacity, CP (J/ g-at K)

Heat capacity, CP (J/ g-at K)

40

Gd5Si2Ge2 M

200 160 120

H=0 Hμ0=0 H = 10kOe T Hμ0=100

80 40 0

0

50

100

150

200

250

300

350

Temperature, T (K)

0

50

100

150

200

Temperature, mperature, T (K)

Fig. 7. Low temperature heat capacities from 5 to 350 K at zero and 100 kOe: (a) Sm5S4; (b) Sm5Si2Ge2; (c) Sm5Ge4; (d) Gd5Si2Ge2.

a magnetic field of 65 kOe applied along the c axis can trigger the PM to FM transition, but along the hard magnetic direction (the b axis) fields >70 kOe are necessary to cause FM ordering in single crystal Tb5(Si2.2Ge1.8) [33]. But no such anisotropy, or large magnetic fields, are required in the Gd5(Si4xGex) single crystals for the formation of the FM state [33]. This difference is due to the strong single ion anisotropy of the 4f wave function in Tb, but is absent in Gd because it has a spherical 4f wave function (i.e. J = S = 7/2, L = 0). The last major difference between the Tb and Gd R5(Si2Ge2) compounds is that the magnetic and structural transitions in the Tb compound are decoupled with an 8 K separation: the magnetic (PM to FM) transition occurs at 111 K and the structural (M to OI) transition occurs at 93 K [31,34]. However, under pressure the two transitions are coupled at 9 kbar and 115 K [34]. 3.3. Er5Si4 At room temperature, only the two orthorhombic phases exist, OII from x ffi 1 to x ffi 4, and OI from x = 0 to x = 0.7,

while the M phase is not observed in the Er5(Si4xGex) system. Although the OII phase dominates in this pseudo-binary system, the most exciting and exotic phenomena are found in Er5Si4, and most of the properties are in direct opposition to what is observed in the Gd5T4 and Tb5T4 systems. When OI Er4Si4 is cooled below room temperature, it undergoes a first-order transformation to the M polymorph at 220 K [35,36] and, upon further cooling, the PM M phase becomes FM via a second-order transformation at 30 K [35]. In this case, the magnetic and structural transformations are uncoupled by 200 K, which is in great contrast to Gd5(Si2Ge2) in which the two transformations are coupled at 270 K (Section 2), and Tb5(Si2Ge2), where they are decoupled by only 8 K (Section 3.2). The 220 K structural transformation which is between two PM polymorphs exhibits an unprecedented magnetic field dependence. Between 0 and 40 kOe, the magnetic field has no effect on the crystallographic transformation, which is normal for a paramagnet, but when the magnetic field is >40 kOe, the transition temperature is lowered from 218.5 K to 215 K at 100 kOe [37]. The anomalous behavior is not fully understood.

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The low temperature 30 K magnetic transition also exhibits unique behavior under an applied magnetic field. In this case, when a magnetic field is 80 kOe or greater, the FM M structure changes to FM OI, which is the zero magnetic field ground state for most of the R5T4 materials studied to date. The existence of a magnetic field induced crystallographic phase change without a change in the spin structure in a magnetic substance is indeed a rare event, and it is the first time it has been observed in the R5T4 alloys. The final unusual property of Er5Si4 is that the magnetic moments are aligned parallel to the b axis in the individual slabs, while in the Gd5T4 and Tb5T4 systems the magnetic moments lie in the ac plane, perpendicular to the b axis [38]. In the low field M phase, the moments are slightly canted away from the b axis, while in the high field OI phase, they are almost perfectly parallel to the b axis. 3.4. Yb5(Si4xGex) The Yb5(Si4xGex) pseudo-binary system, as one might expect from Yb’s multivalent behavior (2+, 3+ and mixed valence) [39], is different from all the other R5T4 systems. All the compositions between and including the two end members, Yb5Si4 and Yb4Ge4 have the same OI crystal structure. This suggests that Yb has a valence less than three, otherwise, if Yb was trivalent, most of the compositions would be expected to have at least the OII structure and perhaps also the M structure, because the Er5T4 has both these structures, and Lu5T4 has only the OII structure, and these two lanthanides lie on either side of Yb. This is evident in Fig. 5a and b, where the lattice parameters of the Yb5T4 two end members in general lie above the curve established by the trivalent lanthanides. From a crystallographic and magnetic susceptibility study, Ahn et al. [40] found that all the alloys order antiferromagnetically at 3 K and that they are heterogeneous mixed valent systems in which 60% of the Yb atoms are divalent and 40% are trivalent. This is another anomaly compared with the other OI R5T4 compounds in that, generally, the latter almost always have a FM ground state. Heat capacity measurements as a function of magnetic field indicate that there is a magnetic field induced FM-like state when the magnetic field exceeds 13 kOe. In a Mo¨ssbauer study of these Yb5(Si4xGex) alloys, Voyer et al. [41] report that Yb2+ and Yb3+ are present in about equal amounts (52 and 48%, respectively), which is slightly different from the percentage estimated from magnetic measurements, as noted above. They also note that the magnetic moments on the Yb3+ ions at 0 K is 2.1lB, and that magnetic ordering develops <1.7 K. 4. Ductile intermetallic compounds 4.1. The discovery of the ductile RM CsCl-type compounds The discovery of ductile intermetallic compounds was also the outgrowth of an applied research project on devel-

25

oping low temperature, <15 K, cryocooler regenerator alloys. Most of the regenerator beds used today consist of spherical powders, but theoretical modeling shows that, if one has parallel plates instead of powders, there is a significant increase in the efficiency/cooling power of the regenerator. One of the important properties of regenerators is that they have a large volumetric heat capacity <20 K, but for most materials the heat capacity drops-off below 50 K, rapidly approaching zero. The only materials which have large heat capacities at or below 620 K are those which have magnetic transitions in the 1–20 K temperature regime, and a number of lanthanide intermetallics are currently being used in cryocoolers to reach temperatures as low as 5 K [42]. But these compounds (e.g. Er3Ni and HoCu2) are brittle, and can only be used as spheres (0.3 mm dia.). Other geometries, such as thin sheets in parallel plates or jelly rolls, or wires or screens, etc., are more efficient, but these forms are difficult to fabricate from a brittle intermetallic compound. With these criteria in mind, work began with Er-based intermetallics because of their high heat capacities and low magnetic ordering temperatures, in the hope of finding a somewhat ductile intermetallic. Based on their known magnetic ordering temperatures, the following materials Er(Cu0.5Ag0.5), were examined: (Ho0.75Er0.25)In3, (Ho0.75Er0.25)Cu2, Er(Ni0.9Co0.1) and Er(Ni0.75Co0.25). Dr A.O. Pecharsky, who prepared the samples, told me on 28 January 2000 that, of the five samples when cut using a low speed diamond saw, Er(Cu0.5Ag0.5) seemed to be somewhat ductile, and the other four were quite brittle. Based on these observations, on 25 February 2000, I took an arc-melted button of ErCu, set it on an anvil and struck it about ten times with a hammer. To my amazement, the sample did not shatter as a normal intermetallic compound would do, but it was flattened by 25%, and only a few visible cracks appeared (see Fig. 8). In order to obtain more ¨ nal made quantitative mechanical property data, Dr O. U some compressive stress–strain measurements on ErCu and reported the values (in MPa) of 78, 187 and 402 for the proportional limit, the 0.2% offset yield point, and the true strength at 10% true strain, respectively, indicating that ErCu was indeed quite ductile. Both Er(Cu0.5Ag0.5) and ErCu have the B2 CsCl-type structure. Because this is one of the most common structures formed by rare earth intermetallic compounds (150 B2 RM compounds are known), there may be many more ductile RM intermetallic compounds. Realizing that this discovery opens a new vista of materials science, and that newly acquired knowledge will provide a major advance in the fundamental understanding of the deformation behavior of intermetallic compounds, a study was undertaken of the stress–strain and fracture toughness behaviors of several RM compounds – YCu, YAg and DyCu [43] (Table 1). Polycrystalline YCu was found to be 25% stronger than YAg, but its ductility was significantly less: 27% vs 11% tensile elongation [43]. In subsequent tests on other polycrystalline YAg samples, it was

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ture toughness, and an intermediate strength, while DyCu has both the highest tensile strength and fracture toughness and an intermediate ductility. For comparison the corresponding values for NiAl and commercial aluminum alloys (e.g. 3105) are also presented in Table 1. The fracture toughness value for the RM compounds is two to five times larger than that of NiAl (which also has the B2 structure) but about the same as that of aluminum alloys. Also, the tensile strength and elongation values for the RM phases and commercial grade aluminum alloys are comparable, but the elongations are five to ten times larger than that of the brittle NiAl intermetallic compound. Fig. 8. ErCu ingot after being struck with 10 hammer blows (4).

4.2. YAg  ‘‘mithril” found that the ductilities range from 17% to 21% elongation, and tensile strengths from 150 to 170 MPa [44]. Studies of YAg [44], YCu [45] and DyCu [46] single crystals, which were elongated 5%, revealed that the primary active slip plane is the {1 1 0} and that the {1 0 0} is a secondary slip plane. Slip lines observed on the surface of the tensile specimens indicate that the slip direction is h0 1 0i. In the case of DyCu, the slip trace analysis also revealed that the {1 0 0}h1 1 1iwas the secondary slip system. Transmission electron microscopy (TEM) examination of the deformed YAg single crystal using a g  b = 0 out of contrast analysis showed that the predominant Burgers vector for the dislocations present in the YAg sample was h1 1 1iwith a minor number of h0 1 1i dislocations [43,44]. However, for DyCu the TEM study revealed both h1 0 0i and h1 1 1i dislocations and that the h1 0 0i was more predominant [46]. Thus, it would appear that the deformation mechanisms are somewhat different in YAg compared with that in DyCu. Fracture toughness studies of polycrystalline YAg, YCu and DyCu along with standard tensile tests showed that there are significant differences between the three RM B2 intermetallic phases [47] (see Table 1). YAg is the most ductile and weakest of the three and has intermediate fracture toughness. YCu is the least ductile and has the lowest frac-

On 1 April 2004, the American Chemical Society named ‘‘mithril” from the Lord of the Rings the ‘‘Molecule of the Week” because an anonymous web-poster (Olog-hai) suggested that YAg might be the material in the real world which comes closest to the magical metal ‘‘mithril” (which is Elvish for ‘‘true silver”) [53]. A few months later, the Boston Museum of Science requested from the Ames Laboratory a coat of mail made of YAg for their ‘‘Lord of the Rings Trilogy” exhibit (1 August–24 October 2004). But because of the prohibitive cost of such a coat the Laboratory suggested that some YAg rings could be made for exhibit. These rings (Fig. 9) were machined from an arcmelted ingot 2.5 cm in diameter, 3 cm long, using a standard lathe. The YAg alloy machined just like brass and work hardened as the center of the ingot was bored out. 4.3. Direct observation of dislocations Although h1 0 0i, h1 1 0i and h1 1 1idislocations had been identified earlier in the RM phases via g  b = 0 analysis (see Section 4.2), the first direct observation of a dislocation in an RM phase was not made until late 2007, using

Table 1 Tensile strength, elongation and fracture toughness of some RM B2 phases Compound or alloy

Tensile strength (MPa)

Elongation (%)

KIC value pffiffiffiffi (MPa m)

YCua DyCub YAgb NiAl Commercial Al alloys

185 205 150 330c 170e

11 16 23 25c 20e

12.0 25.5 ± 2.3 19.1 ± 0.2 5.8 ± 0.7d 20–45f

a YCu is polymorphic, it transforms to the orthorhombic B27 FeB-type structure at 140 K. b DyCu and YAg are monomorphic. c After Hahn and Vedula [49]. d After Kim et al. [50]. e Metals Handbook [51]. f Metals Handbook [52].

Fig. 9. YAg (‘‘mithril”) rings, 22 mm o.d.  20 mm i.d. for the Boston Museum of Science’s ‘‘Lord of the Rings Trilogy” exhibit (1 August–24 October 2004).

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high-resolution transmission electron microscopy (TEM) [54]. Xie and co-workers observed that superlattice dislocations in YAg had dissociated into two ½h1 1 0i superpartials bound to either a stacking fault or an antiphase boundary (APB) in a (1 0 0) plane. Assuming that YAg is isotropic, the ½h1 1 0i{1 1 0} stacking fault energy is 287 mJ m2 for pure screw partials or 422 mJ m2 for pure edge partials. These values are in good agreement with those calculated from first principles: 315 mJ m2 (see Section 4.5). Thus, it was concluded that the stacking faults and antiphase boundaries are major contributors to the large strain-hardening rate in YAg. However, other plastic deformation mechanisms, such as deformation twining may also play a role. 4.4. Brittle RM phases Most of the RM phases studied as of the end of 2007 were found to be ductile (16), and a few have been found to be brittle, namely YMg, CeMg, YZn, (Tb0.88Dy0.12)Zn, and ScRu. Russell et al. [45] reported that all eight of the (Tb0.88Dy0.12)Zn single crystals studied failed by brittle fracture without any noticeable elongation in tensile tests. Three different orientations [1 0 0], [2 1 1] and [1 1 1] were examined. The tensile strength along the [1 0 0] was the lowest 115 MPa, and that in the [2 1 1] direction was the largest 215 MPa, while in the [1 1 1] the tensile strength was somewhat less than that in the [2 1 1], 180 MPa. No slip lines were observed on the specimen faces, and all specimens had cleavaged fracture surfaces. The authors also ruled out the possibility of hydrogen embrittlement in the RZn alloy because the hydrogen content (40 wt.ppm) was about the same as in the YCu single crystals which had ductilities of 7–10%. The reason for brittle failure in some RM B2 phases and ductility in other RM CsCl-type B2 phases is not known but, as more experimental data are obtained, trends should develop and give us some clues. In addition, first-principles calculations will also be quite useful in this regard (see Section 4.5). 4.5. Ab initio calculations The results of first-principles calculations on YCu and YAg were reported by Morris et al. [55]. These authors calculated: the lattice parameters; elastic constants; the relative stability of the B2 CsCl-type structure with respect to the B27 FeB-type structure; and the APB and unstable stacking fault energies. The calculated lattice parameters and elastic constants were in good to excellent agreement with the experimental values; the largest discrepancy was for the c12 value of YAg, which was 8% lower than the experimental value. The 0 K energies of the B2 and B27 structures of YCu and YAg were calculated and that of the B27 phase for YCu is lower than that of the B2 phase by 0.038 eV per formula unit, which is consistent with the fact that YCu exhib-

27

its a B2 ? B27 phase transition at 140 K [48]. However, YAg is monomorphic and only crystallizes in the CsCl-type structure, and the calculated energies are consistent with this fact; the B2 structure has a lower energy than that of the B27 structure by 0.20 eV. The APB and unstable stacking fault energies were also calculated for two of the ductile RM phases (YCu and YAg) and the brittle NiAl, which also has the B2 CsCl-type structure. The APB energies for the 12 h1 1 1i{1 1 0} and 12 h1 1 1i{1 1 2} partial dislocations were comparable for the three compounds, ranging from 745 (for YAg) to 1030 mJ m2 (for YCu), and from 680 (for YAg) to 1090 mJ m2 (for YCu), respectively, with the NiAl values lying between the YM phase values. Thus, the APB energies cannot account for the observed ductile/brittle characteristics of these three phases. In contrast, however, there is a clear distinction between the unstable stacking fault energies of the ductile YCu and YAg phases and the brittle NiAl. For the 12 h1 0 0i{0 1 0} stacking fault, the values are 700, 560 and 1835 mJ m2, respectively, and for the 12 h1 0 0i{0 1 1} the values are 325, 315 and 1290 mJ m2, respectively, a fourfold decrease for the ductile RM intermetallics compared with NiAl. Theoretical calculations have been a valuable tool in helping to understand the ductile nature of these B2 RM phases. Continued strong interactions between the theorists and experimentalists are critical in reaching a more complete knowledge of the extraordinary mechanical behavior of these rare earth ductile intermetallics, and this should be a valuable first step in our understanding of the nature of failure of materials and the limitations in using intermetallic compounds in structural and other applications. 5. Summary and conclusions Over the years, in response to other scientists from private industry requesting the design of rare earth materials with specific magnetic behaviors, my colleagues at the Ames Laboratory and I were able to provide them with rare earth alloys which met their specifications. In discovering, testing and developing these materials for a particular application, we sometimes observed unusual behaviors, which at the time were not understood. In order to remedy this lack of knowledge, basic research investigations were carried out. This led to important new discoveries, such as the GMCE in 1997 and ductile intermetallic compounds in 2003. Concomitant with the former, an applied research project showed that magnetic refrigeration is a viable cooling technology and may soon become competitive with current compressed gas refrigeration and cooling processes. Acknowledgements These accomplishments would not have been possible without the financial support of the US Department of Energy (primarily through the Office of Basic Energy Sciences, Materials Sciences Division, and its predecessor

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organizational units) who have funded me and my group from my beginning graduate student days until today (a total of 56 years). Currently, our research is supported by the Office of Basic Energy Sciences, Materials Sciences Division of the United States Department of Energy under contract No. DE-AC02–07CH11358 with Iowa State University. I wish to thank my close colleagues Vitalij Pecharsky and Alan Russell for their comments and our interesting discussions on our ongoing research projects, some of which are discussed above. I also acknowledge the assistance of Yaroslav Mudryk in preparing some of the figures in this paper. I wish to express my gratitude to several organizations: Acta Materialia, Inc., for choosing me as the recipient of the 2008 Acta Materialia Gold Medal; Elsevier for minting and providing the medal; and TMS for providing the venue for receiving this recognition and the special symposium at the annual meeting in New Orleans, 9–12 March 2008. The stellar efforts of the symposium organizers, Vitalij K. Pecharsky, Jamie R. Morris and Ashutosh Tiwari, are deeply appreciated. I should also acknowledge the faith, trust and efforts of my nominator for the Acta Gold Medal, T.B. Massalski. Finally my wife, Melba, and my children, who are now deeply committed to their respective careers, Tom, Dave, Ed and Kathy, also deserve a great deal of credit for their patience, understanding and support. References [1] Gschneidner Jr KA. J. Alloys Compd 1992;180:1. [2] Gschneidner Jr KA, Pecharsky VK. Intern J Refrig 2008;31:945. [3] Zimm C, Jastrab A, Sternberg A, Pecharsky V, Gschneidner Jr KA, Osborn M, et al. Adv Cryog Eng 1998;43:1759. [4] Pecharsky VK, Gschneidner Jr KA. Phys Rev Lett 1997;78:4487. [5] Pecharsky VK, Gschneidner Jr KA. Appl Phys Lett 1997;70:3299. [6] Holtzberg F, Gambino RJ, McGuire TR. J Phys Chem Solids 1967;28:2283. [7] Pecharsky VK, Gschneidner Jr KA. J Alloys Compd 1997;260:98. [8] Choe W, Pecharsky VK, Pecharsky AO, Gschneidner Jr KA, Young Jr VG, et al. Phys Rev Lett 2000;84:4617. [9] Pecharsky VK, Gschneidner Jr KA. Pure Appl Chem 2007;79:1383. [10] Morellon L, Algarabel PA, Ibarra MR, Blasco J, Garcia-Landa B, Arnold Z, et al. Phys Rev B 1998;58:R14721. [11] Morellon L, Blasco J, Algarabel PA, Ibarra MR. Phys Rev B 2000;62:1022. [12] Pecharsky VK, Gschneidner Jr KA. Adv Mater 2001;13:683. [13] Morellon L, Stankiewicz J, Garcia-Landa B, Algarabel PA, Ibarra MR. Appl Phys Lett 1998;73:3462. [14] Levin EM, Pecharsky VK, Gschneidner Jr KA. Phys Rev B 1999;60:7993. [15] Levin EM, Pecharsky VK, Gschneidner Jr KA, Tomlinson P. J Magn Magn Mater 2000;210:181. [16] Day C. Phys Today 2007;60(12):12. [17] Levin EM, Pecharsky VK, Gschneidner Jr KA. Phys Rev B 2001;63:174110. [18] Zou M, Tang H, Schlagel DL, Lograsso TA, Gschneidner Jr KA, Pecharsky VK. J Appl Phys 2006;99:08B304. [19] Haskel D, Lee YB, Harmon BN, Islam Z, Lang JC, Srajer G, et al. Phys Rev Lett 2007;98:247205.

[20] Paudyal D, Pecharsky VK, Gschneidner Jr KA, Harmon BN. Phys Rev B 2007;75:094427. [21] Tan L, Kreyssig A, Kim JW, Goldman AI, McQueeney RJ, Wermeille D, et al. Phys Rev B 2005;71:214408. [22] Roy SB, Chattopadhyay MK, Chaddah P, Moore JD, Perkins GK, Cohen LF, et al. Phys Rev B 2006;74:012403. [23] Roy SB, Chattopadhyay MK, Banerjee A, Chaddah P, Moore JD, Perkins GK, et al. Phys Rev B 2007;75:184410. [24] Pecharsky VK, Holm AP, Gschneidner Jr KA, Rink R. Phys Rev Lett 2003;91:197204. [25] Ouyang ZW, Pecharsky VK, Gschneidner Jr KA, Schlagel DL, Lograsso TA. Phys Rev B 2006;74:094404. [26] Griffiths RB. Phys Rev Lett 1969;23:17. [27] Ugurlu O, Chumbley LS, Schlagel DL, Lograsso T. Acta Mater 2006;54:1271. [28] Smith GS, Thorp AG, Johnson Q. Acta Cryst 1967;22:940. [29] Ahn K, Pecharsky VK, Gschneidner Jr KA. Phys Rev B 2007;76:014415. [30] de Gennes PG. Compt Rend 1958;247:1836. [31] Morellon L, Ritter C, Magen C, Algarabel PA, Ibarra MR. Phys Rev B 2003;68:024417. [32] Garlea VO, Zarestky JL, Jones CY, Lin L-L, Schlagel DL, Lograsso TA, et al. Phys Rev B 2005;72:104431. [33] Zou M, Mudryk Ya, Pecharsky VK, Gschneidner Jr KA, Schlagel DL, Lograsso TA. Phys Rev B 2007;75:024418. [34] Morellon L, Arnold Z, Magen C, Ritter C, Prokhnenko O, Skorokhod Y, et al. Kamarad J Phys Rev Lett 2004;93:137201. [35] Pecharsky VK, Pecharsky AO, Mozharivskyj Y, Gschneidner Jr KA, Miller GJ. Phys Rev Lett 2003;91:207205. [36] Mozharivskyj Y, Pecharsky AO, Pecharsky VK, Miller GJ, Gschneidner Jr KA. Phys Rev B 2004;69:144102. [37] Pecharsky AO, Gschneidner Jr KA, Pecharsky VK, Schlagel DL, Lograsso TA. Phys Rev B 2004;70:144419. [38] Magen C, Ritter C, Morellon L, Algarabel PA, Ibarra MR, Tsokol AO, et al. Phys Rev B 2006;74:174413. [39] Gschneidner Jr KA. J Less-Comm Metals 1969;17:13. [40] Ahn K, Tsokol AO, Mozharivskyj Y, Gschneidner Jr KA, Pecharsky VK. Phys Rev B 2005;72:054404. [41] Voyer CJ, Ryan DH, Ahn K, Gschneidner Jr KA, Pecharsky VK. Phys Rev B 2006;73:174422. [42] Gschneidner Jr KA, Pecharsky AO, Pecharsky VK. In: Ross Jr RG, editor. Cryocoolers 12. New York: Kluwer Academic/Plenum; 2003. p. 457. [43] Gschneidner Jr KA, Russell A, Pecharsky A, Morris J, Zhang Z, Lograsso T, et al. Nature Mat 2003;2:587. [44] Russell AM, Zhang Z, Lograsso TA, Lo CCH, Pecharsky AO, Morris JR, et al. Acta Mater 2004;52:4033. [45] Russell AM, Zhang Z, Gschneidner Jr KA, Lograsso TA, Pecharsky AO, Slager AJ, et al. Intermetallics 2005;13:565. [46] Cao GH, Shechtman D, Wu DM, Becker AT, Chumbley LS, Lograsso TA, et al. Acta Mater 2007;55:3765. [47] Zhang Z, Russell AM, Biner SB, Gschneidner Jr KA, CCH Lo. Intermetallics 2005;13:559. [48] Ritter C, Ibarra MR, Ibberson RM. J Phys: Condens Mat 1992;4:L39. [49] Hahn KH, Vedula K. Scripta Metall 1989;23:7. [50] Kim T, Hong KT, Lee KS. Intermetallics 2003;11:33. [51] Metals handbook, vol 2. 10th ed. Materials Park, OH: ASM International; 1990. p. 87. [52] Metals handbook, vol 19. 10th ed. Materials Park, OH: ASM International; 1996. p. 774. [53] American Chemical Society named YAg the ‘‘Molecule of the Week – Mithril” on April 1, 2004; also see http://greenbooks.theonering.net/ guest/files/103003_02.html. [54] Xie S, Russell AM, Becker AT, Gschneidner Jr KA. Scripta Mater 2008;58:1066. [55] Morris JR, Ye Y, Lee YB, Harmon BN, Gschneidner Jr KA, Russell AM. Acta Mater 2004;52:4849.