Ramanathan Sh. (Ed.) Thin Film Metal-Oxides: Fundamentals and Applications in Electronics and Energy

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7 High-T

Superconducting Thin- and Thick-Film–Based Coated Conductors 241

7.4 Enhancement of Self-Field J

If the objective is to maximize performance of HTS wires in transmission lines,

which operate under self ﬁelds up to 0.1 T, we need to optimize the ﬁrst regime in

the J

veesus H curve and increase J

.H D 0/ as close as possible to the depairing

current density. This means that the HTS ﬁlm needs to be highly crystalline, dense,

and precipitate-free. The ﬁlm should resemble a single crystal with exception of

a few, largely spaced (1; 000

A apart) extended defects along c that act as deep

and narrow potential wells, strongly pinning the ﬂux lines but taking away very

little superconducting material from the nearly perfect surrounding matrix. One of

the ﬁrst studies to illustrate the importance of c-axis columnar defects produced

by threading dislocations in the low ﬁeld regime was that of Dam et al. [36]. In

this study, the density of threading edge dislocations occurring predominantly at

the boundaries between growing islands was changed by changing the size of the

islands. This was achieved by appropriate tuning of deposition conditions. Although

little variations in J

.H D 0/ were observed, increasing the density of dislocations

extended the J

plateau to higher ﬁelds.

The RBCO growth mode, on SrTiO

(STO) as well as other oxide buffer layers,

involves an initial two-dimensional nucleation and growth that only lasts for the

ﬁrst 5–15 nm, after which a 3D island growth mode takes place. This evolution of

the growth mode was observed on both TiO

- and SrO-terminated STO surfaces.

In addition, as it has been documented for RBCO growth on TiO

-terminated STO

substrates, even at the initial stage of RBCO deposition, formation of CuO

pre-

cipitates occurs. This happens because nucleation of RBCO on TiO

-terminated

STO always starts with a BaCuO perovskite block leading to a termination of

the ﬁrst unit cell with a CuO or CuO

layer. Because the chain layer is unstable

at the oxygen partial pressures normally used for growth, the system lowers its

free energy by precipitating Cu oxide islands on a continuous BaO layer, there-

fore disrupting layer-by-layer growth [37, 38]. The tendency to abandon a 2D

growth mode in favor of islands formation seems to be unavoidable in HTS in-

dependently on substrate surface and growth conditions. Nevertheless, the atomic

smoothness of the ﬁrst 10 unit cells of RBCO can be greatly improved by

careful composition control during deposition of the ﬁrst atomic layers. Rijnders

et al. [39] have systematically investigated the stacking sequence of RBCO per-

ovskite blocks on STO and experimentally determined the following sequences.

On TiO

-terminated STO, two stacking sequences of atomic layers at the interface

are possible: bulk–SrO–TiO

–BaO–CuO–BaO–CuO

–R–CuO

–BaO–bulk,re-

ferred as R133, and bulk–SrO–TiO

–BaO–CuO

–R–CuO

–BaO–bulk, referred

as R122. On SrO-terminated STO two different stacking sequences are in-

stead likely: bulk–SrO–TiO

–SrO–CuO–BaO–CuO

–R–CuO

–BaO–bulk, and

SrO–TiO

–SrO–CuO

–R–CuO

–BaO–CuO–BaO–bulk. On as-received STO sub-

strates both SrO and TiO

termination of the surface are present and four possible

stacking sequences at the interface (SSI) occur. Multiple stacking sequences lead

to increase roughness and antiphase boundaries, which develop when neighboring

islands with different SSI coalesce. A way to reduce the occurrence of APBs and

242 C. Cantoni and A. Goyal

promote 2D growth is to use a single surface termination, TiO

, and “impose”

the stacking sequence of the ﬁrst RBCO layer. The ﬁrst unit cell RBCO layer is

deposited using a cation ratio RWBaWCu D 1W2W2, forcing the SSI to be R122, above

which regular 123 stoichiometry RBCO is deposited. Atomic layer control of the

growth process is achieved by monitoring in situ the intensity oscillations of diffrac-

tion spots obtained through reﬂection of high-energy electron diffraction (RHEED).

RBCO ﬁlms grown using this procedure are signiﬁcantly smoother and denser

than ordinary ﬁlms but their superconducting properties are nonetheless worse.

In particular, T

is suppressed due to the difﬁculty in achieving the right oxygen

stoichiometry in such dense ﬁlms. Moreover, as reported by Haage et al. [40]and

Cantoni et al. [41], interfaces formed by APBs are strong pinning sites for vortices

and their presence can greatly enhance J

at very low magnetic ﬁelds. Therefore,

suppression of APBs by controlling the growth of the ﬁrst RBCO monolayers is not

likely to produce HTS ﬁlms that show high density and crystal quality, and high J

at the same time.

Another approach that has recently been used to obtain dense YBCO ﬁlms with

only a few very strong extended defects per m

is to take advantage of the sig-

niﬁcant lattice strain that develops in ﬁlms grown on vicinal surfaces produced by

a substrate miscut angles of a few degrees. Wu et al. [42] have shown that after a

thickness of about 5–10 nm, when the APBs originated in the YBCO by the substrate

steps have been largely annihilated by stacking faults, nanometer size precipitates

and, subsequently, nano-pores develop in the growing ﬁlm with a density of 3–8

pores/m

. While nanopores occurring in YBCO ﬁlms on ﬂat (non-vicinal) sub-

strates annihilate after a thickness of a few hundred nanometer, pores in YBCO ﬁlms

grown on vicinal surfaces extend for the entire ﬁlm thickness up to a few microns.

Both nanoparticles nucleation and pore evolution are expected to be intimately con-

nected to the strain generated in the ﬁlm during the growth process on the vicinal

surface. The nanopores act as strong and large columnar defect providing strong

ﬂux pinning according to a mechanism investigated in highly crystalline ultrathin

lead ﬁlms with nanovoids [43]. At the same time, the low density of pores results

in both a large superconducting cross-section and a high T

. As a consequence, J

values in self-ﬁeld measured for these ﬁlm are as high as 8:3 MA=cm

, signiﬁcantly

higher than typical self-ﬁeld J

sof4–5 MA=cm

on control YBCO samples, and

higher than the best J

ever reported on undoped YBCO ﬁlms [44]. An image of

nanopores produced on a 10

miscut STO substrate as obtained by atomic force

microscopy (AFM) is shown in Fig. 7.4.

All the high-J

HTS ﬁlms described above have a thickness in the range

0:1–0:3 m, which is considered optimal for achieving the largest J

s. As the

superconductor thickness increases above 1m, the critical current density de-

creases due to both extrinsic and intrinsic material issues. However, for practical

application of HTS, it is essential to increase the overall current carrying capacity

of the coated conductor to meet speciﬁc requirements. The ﬁgure of performance

of a coated conductor is not the J

of the RBCO ﬁlm but the engineering current

density J

of the whole conductor, which is the total current carried by the super-

conductor, I

, divided by the total cross-section of the conductor. The thickness

7 High-T

Superconducting Thin- and Thick-Film–Based Coated Conductors 243

Fig. 7.4 AFM image of a 200-nm-thick porous YBCO ﬁlm on a 10

vicinal STO substrate in-

cluding the ﬁlm’s vicinal direction. Its inset shows a TEM image of the ﬁlm/substrate interface of

a porous YBCO ﬁlm on a 15

STO substrate. (a) AFM images on (b) an 8-nm thick and (c) a 16-

nm-thick YBCO ﬁlm on a 20

vicinal STO substrate, respectively. The nanoparticles are observed

on (b) and pores just initiated atop nanoparticles are circled on (c). The scales for (b)and(c)are

the same. The insets show the AFM image (bottom) and depth proﬁle (top) on a selected pore. The

x-direction in the top inset is the scan direction deﬁned by the green line in the bottom inset. A

nanoparticle at the bottom of the pore can be clearly seen. Reprinted with permission from J. Z.

of the superconducting ﬁlm is only a small fraction (2–3%) of the total conductor

thickness, including the substrate and the stabilizing metal overlayer (see Fig. 7.1),

and it is not practical or even possible for every ﬁlm deposition method to make a

superconducting coating more than 5m thick. Therefore, the only option left for

increasing the J

of the entire conductor is to produce HTS ﬁlms with a thickness

of a few microns that retain most of the high J

of thinner ﬁlms. As of today, the

origin of the J

decrease with thickness is not completely understood. Initially,

this effect has been attributed to microstructural evolution of the YBCO ﬁlms with

increasing thickness and resulting roughening of the ﬁlm/vapor interface, coarsen-

ing of precipitates, secondary outgrowths, and crystallinity degradation. However,

even when these material issues are largely mitigated, J

still decreases rapidly

with thickness. Recently, (Y, Eu)BCO ﬁlms have shown a self-ﬁeld J

thickness

dependence quantitatively similar to that of YBCO and, at the same time, a much

denser and smoother microstructure with greatly reduced surface roughness, and

larger and well-connected ﬁlm islands [45]. However, these observations might still

reconcile with the explanation of the J

decrease with thickness that attributes the

suppression of J

to the intrinsic healing with ﬁlm thickness of threading misﬁt

dislocations and APBs originating at the substrate interface. As thickness increases,

the REBCO ﬁlm exhibits fewer ﬂux pinning defects and J

decreases. For this

reason, it is important to introduce an additional, artiﬁcial pinning mechanism in

the growing ﬁlm that can be retained at large enough thicknesses. Substituting an

RE element for Y or using a mixture of rare earths instead of Y are examples of

one way in which artiﬁcial pinning can be introduced. In this case, the ion size

variation introduces local strain ﬁelds in the ﬁlm lattice that might be responsible

244 C. Cantoni and A. Goyal

for additional random pinning, yielding better performance. The additional pinning

in (Y, Eu)BCO ﬁlms might explain why the critical current thickness dependence

in these ﬁlms remains comparable to that of YBCO even if larger island size and

increased lattice parameter as compared to YBCO should lead to lower interfacial

critical current enhancement.

With exception of transmission cables, all the other power applications of HTS

wires, such as motors and generators, require operation in intermediate magnetic

ﬁelds. For these applications, it is important to enhance pinning in a range of ap-

plied magnetic ﬁeld of 3–5T, which requires introduction of much denser arrays of

artiﬁcial pinning sites with consequent unavoidable decrease of T

and supercon-

ducting cross-section, and consequently self-ﬁeld J

7.5 Introduction of Artiﬁcial Flux Pinning Nanostructures

for Enhancement of In-Field J

When substantial magnetic ﬁelds are involved, the direction of the ﬁeld with respect

to the basal plane of the ﬁlm acquires relevance. Because of the anisotropic crystal

structure of HTS, J

in as-grown YBCO ﬁlms is maximal when the magnetic ﬁeld

is directed parallel to the ab planes, a situation which is not preferred in all power

applications. Figure 7.5 shows a typical curve for J

as function of magnetic ﬁeld

Fig. 7.5 Angular dependence of J

for a 1:55-m-thick YBCO ﬁlm on SrTiO

at 5 T. Note that the

largest current density is obtained when the ﬁeld lies in the a,b plane. The solid line represents the

random pinning contribution. Reprinted with permission from L. Civale et al. Physica C 412–414

used in the measurements

7 High-T

Superconducting Thin- and Thick-Film–Based Coated Conductors 245

orientation for a YBCO ﬁlm grown by pulsed laser deposition (PLD) on SrTiO

Superimposed to the data points is a curve that represents the contribution to J

from uncorrelated pinning, produced by randomly distributed localized defects. The

random pinning contribution is obtained by ﬁtting the J

.™/ data to the scaling curve

predicted for electronic mass anisotropy effects alone, according to the analysis of

Civale et al. [46]. The difference in J

between the data and the random pinning

component is due to correlated pinning provided by extended, parallel pinning cen-

ters, either linear or planar, which naturally arise during growth. These defects are

usually oriented both along the c-axis (especially for PLD ﬁlms) and parallel to the

ab planes. A ﬁnal contribution to the peak at H//ab is given by intrinsic pinning,

also due to the layered crystal structure.

In view of meeting power applications criteria, today’s research goal is to intro-

duce artiﬁcial pinning sites that enhance J

for orientations close to H//c near to the

level of J

at H//ab, giving rise to a more isotropic angular dependence.

A successful approach for increase pinning in magnetic ﬁeld is to introduce non-

superconducting, chemically compatible particles of appropriate size throughout

the superconductor matrix. This can be accomplished by deposition from multi-

ple sources, in the case of in situ methods, or by adding excess elements to the

superconductor precursor, in the case of ex situ methods. In situ laser ablation

from a sintered superconductor target already containing the nanoparticle material

in a disperse form is another option. While uncorrelated pinning from randomly

distributed point-like defects is generally weak, randomly dispersed, ﬁnite size par-

ticles can provide strong pinning. Several examples of impurity additions in the

form of nanoparticles and/or nanodots are available. In one approach, a YBCO non-

superconducting phase Y

BaCuO

, also known as 211, was added by intermittently

depositing YBCO layers and discontinuous 211 layers by PLD using two deposition

targets [47, 48]. This technique resulted in 211 nanoparticles with an average size

of 15 nm distributed in the YBCO ﬁlm, which produced little or no enhancement in

pinning along c but signiﬁcant improvements for H//ab. Correlated pinning along

the c-axis was observed when randomly distributed BaZrO

(BZO) particles with a

modal size of 10 nm were inserted in the YBCO ﬁlm by ablating from a PLD tar-

get with Zr oxide additions [49]. In the last study, the observed correlated pinning

along the c-axis was related to misﬁt dislocations between the nanoparticles and the

superconducting matrix, which aligned along the c-direction, with an areal density

of 400 m

2

compared to 80 m

2

for undoped YBCO ﬁlms.

7.5.1 Self-Assembled Nanostructures

When BZO nanoparticles have an epitaxial relationship to the YBCO matrix,

a signiﬁcant strain ﬁeld results in the growing ﬁlm due to the large mismatch

between the two lattices, which is about 8% when BZO(001)//YBCO(001) and

BZO(100)//YBCO(100). Such high strain ﬁelds can induce a remarkable self-

assembly of the nanodots resulting in the growth of almost continuous and uniform

246 C. Cantoni and A. Goyal

Fig. 7.6 On the left: cross-section TEM image of a YBCO ﬁlm on RABiTS with self-assembled

BaZrO

(BZO) nanodots and nanorods (left). The image shows that the BZO nanodots are aligned

along the c-axis of YBCO and are about 2–3 nm in diameter. On the right: Z-contrast STEM image

of a single BZO nanodot. Four misﬁt dislocations can be seen around the BZO nanodot. The extra

semiplane in the edge dislocation cores are marked in blue. Different sets of planes in the YBCO

and BZO are indicated. The zone axis is parallel to the YBCO c-axis. Adapted from ref. [50]

BZO nanorods in the YBCO matrix. Goyal et al. [50, 51] were able to produce

such an ordered array of BZO nanodots in YBCO via laser ablation from a single

target comprising a mixture of YBCO powder and BZO nanoparticles. During si-

multaneous deposition of YBCO and BZO, the two phases were found to separate

in the growing ﬁlm and interact through the strain ﬁeld similar to consecutively

deposited phases. Similar nanostructures were obtained using different oxides in

addition to BZO (CaZrO

,YSZ,MgO,Ba

1x

TiO

) establishing the general va-

lidity of this approach [52]. Theoretical formulations to explain how self-assembly

occurs in the YBCO-BZO system via both energetic and kinetic arguments have

been developed [53].

Figure 7.6 shows two TEM images obtained from a YBCO/BZO composite ﬁlm

deposited from a YBCO target doped with 2 vol.% BZO nanoparticles. The im-

age on the left is a HRTEM of a cross-section of the ﬁlm, which clearly shows the

BZO columns produced by stacking of nanodots and continuous rods. We note that

the nanocolumns are aligned along the c-axis of YBCO and extend throughout the

thickness of the ﬁlm. The right image in Fig. 7.6 is a high-resolution Z-contrast

STEM image obtained from a plan view specimen of the ﬁlm and shows a single

BZO nanorod section with the c-axis directed perpendicular to the image plane.

Different atomic columns are distinguishable thanks to the dependence of the im-

age intensity on the atomic number. In the YBCO matrix, the Ba/Y columns appear

brighter than the Cu/O columns and within the BZO, BaO columns appear brighter

than ZrO. The image shows that the BZO is epitaxially oriented and coherent with

the YBCO matrix, and misﬁt dislocations form to accommodate mismatch at the

7 High-T

Superconducting Thin- and Thick-Film–Based Coated Conductors 247

H//ab

H//c

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

-40

-20

20 40 60 80 100

120

Undoped

2% BZO

Angle (deg)

H//ab

H//c

H//ab

H//c

Undoped

2% BZO

H//ab

H//c

0.01

0.1

0.01 0.1

Undo ped, Normalized, H//c

2% BZO, No rmalized, H//c

Field (T)

BZO doped

αα

α ~ 0.31

Undoped

αα

α ~ 0.5

Undo ped, Normalized, H//c

2% BZO, No rmalized, H//c

BZO doped

αα

α ~ 0.31

Undoped

αα

α ~ 0.5

Undo ped, No rmalized, H//c

2% BZO, Normalized, H//c

Undoped

α ~ 0.5

(sf)

(MA/cm

)

BZO doped

α ~ 0.31

Fig. 7.7 Left: log–log plot of normalized J

as a function of applied ﬁeld showing that the power

low exponent in the relation J

H

’

is lower in the YBCO ﬁlm with self-assembled nanodots

than for the control, undoped YBCO ﬁlm. Right: angular dependence of J

at 77 K, 1 T for an un-

doped YBCO ﬁlm on RABiTS compared to a YBCO ﬁlm with 2 vol.% BZO nanodots. Signiﬁcant

enhancement of J

is seen at all angles, particularly at H//c (angle D 0

). Both superconducting

ﬁlms have a thickness of 0:2 m. Adapted from ref. [50]

corners of the BZO nanodot. Since the nanodots are aligned along the YBCO

c-direction, the strain from the misﬁt dislocations is also aligned and extended

along this direction. Both the arrays of dislocations and the nanodots themselves

are expected to form ideal ﬂux pinning columns, similar to damage tracks by heavy

ion irradiation [54, 55].

Figure 7.7 shows the pinning enhancement, and compares J

dependence on

ﬁeld magnitude and orientation for a pure YBCO ﬁlm on RABiTS and a BZO-

incorporated YBCO ﬁlm, both grown on the same substrate. The ﬁlms had a

thickness of 0:2 m. It can be clearly seen that the correlated c-axis peak in the

angular dependence of J

is very pronounced for the YBCO ﬁlm containing the

BZO nanodots, indicative of strong pinning defects along the c-axis.

Self-assembly of BZO nanodots can persist for a thickness much larger than

0:2 m, supplying the pinning necessary to compensate for some of the J

decrease

versus thickness and allowing signiﬁcant I

increase. Figure 7.8 shows cross-section

TEM images of a 4m-thick YBCO ﬁlm with incorporated BZO nanorods. Here,

the columns of self-assembled BZO nanodots are not fully aligned along the c-axis

of YBCO but rather show a ‘splayed’ defect microstructure. Such a ‘splayed’ mi-

crostructure is actually more favorable for ﬂux pinning at all ﬁeld orientations [56].

The performance of this 4-m-thick ﬁlm on a short section of CC template (I

353 Acm

1

at T D 65 K;HD 3 T) is higher than record values previously re-

ported and exceeded the minimum requirements for transmission cables, electric

ship propulsion systems, large-scale motors, and generators [57]. The signiﬁcance

of this result is now well apparent with the BZO nanodots technology successfully

transferred from ORNL to SuperPower, one of the leading US company in 2G

248 C. Cantoni and A. Goyal

Fig. 7.8 Lower (a) and higher magniﬁcation (b) cross-section HRTEM images of BZO

nanocolumns in a 4-m-thick YBCO ﬁlm. The columns extend through the entire thickness of

the ﬁlm and show small deviations in their direction from the YBCO c-direction. Adapted from

ref. [57]

coated conductors manufacture [58]. As expected, the self-assembly of BZO dots

is not conﬁned to the PLD process but also occurs, with proper strain-ﬁeld tuning

during YBCO growth by MOCVD, a highly scalable technique used by Superpower

to produce kilometer-long superconducting cables.

While BZO nanocolumns provide strongly enhanced c-axis correlated pinning

with pronounced peak in J

for H//c, little enhancement of pinning for H//ab is usu-

ally observed in any RBCO ﬁlms with columnar defects [59–61]. Some RBCO ﬁlms

even exhibit an “inverse anisotropy” with the highest J

at H//c and the lowest J

H//ab in the temperature ﬁeld regimes of 75–77K, 1 T [59,61]. In order to enhance

for H//ab, controlled introduction of nanoscale defects which are aligned parallel

to the ab-plane is needed. Fortunately, the strain surrounding the BZO precipitates

and responsible for the self-assembly phenomenon can be tuned in such a way that

alignment of the nanodots in the basal plane instead that along c occurs [62]. Films

with such long-range alignment of the BZO dots perpendicular to c show a strong

peak in J

for H//ab. In an effort to combine the two pinning effect and therefore

obtain a more isotropic J

dependence on angle, Wee et al. have shown that it is pos-

sible to fabricate hybrid NdBa

7•

(NdBCO) ﬁlms wherein half of the ﬁlm

has defects structures optimized for H//c and the other half of the ﬁlm has defect

structures optimized for H//ab [63].

The hybrid NdBCO ﬁlms were obtained by in situ deposition of two distinctive

layers each comprising half the thickness of the ﬁlm. The ﬁrst layer was a multilayer

of ŒNdBCO

3 unit cell

=BZO

0:5 unit cell

0:42-m thick and the second layer was a BZO-

doped NdBCO layer 0:43-mthick.

Figure 7.9 shows cross-section TEM images for the hybrid NdBCO ﬁlm in two

different magniﬁcations. BZO columns of 7 10 nm in width are evident in the

top layer, while in the bottom layer BZO nanodot arrays are observed to have the

long-range alignment parallel to the ab-plane. Figure 7.9c illustrates the J

angular

7 High-T

Superconducting Thin- and Thick-Film–Based Coated Conductors 249

Fig. 7.9 Cross-section transmission electron micrographs of 0:85-m-thick NdBCO ﬁlm with a

hybrid nanoscale defect structure. (a) Full cross-sectional area of entire NdBCO layer on IBAD-

MgO templates with conﬁguration of LaMnO

=MgO=Al

/Hastelloy. (b) High magniﬁcation

image of hybrid NdBCO ﬁlm showing a high density of columnar defects of c-axis oriented BZO

nanodots in the top layer of NdBCOCBZO, and a long-range order of BZO nanodot arrays aligned

the ab-plane in the bottom layer of NdBCO/BZO. (c) Angular dependence of J

at 77 K and 1 T

for the hybrid NdBCO ﬁlm in a magnetic ﬁeld applied 30

to 110

with respect to the c-axis

of the ﬁlm. For comparison, the J

data for pure NdBCO and NdBCO

vol.% BZO samples are

also illustrated. Adapted from ref. [63]

dependence for the hybrid ﬁlm as compared to an NdBCO ﬁlm grown in the same

conditions as the top layer (only BZO nanocolumns aligned along c are present),

and a pure NdBCO ﬁlm. All curves are measured by transport at T D 77 Kand

in an applied magnetic ﬁeld of 1 T. The hybrid ﬁlms clearly show a reduced pin-

ning anisotropy, and, consequently a signiﬁcantly higher minimum level of J

, with

respect to all ﬁeld angles.

250 C. Cantoni and A. Goyal

7.6 Concluding Remarks

Based on the results reported in this chapter, it is clear that the development of

coated conductors using the technology of epitaxial ﬁlms of high temperature su-

perconductors has been successful, and that future optimization to meet application

standards looks feasible. In particular, the performance of coated conductors has

been rising steadily as the understanding of the role of ﬂux-pinning defects and

the ability to manipulate them has advanced. However, we are still far from ﬁnal

answers about J

optimization in these materials, especially considering that a the-

oretical model of high-T

superconductivity is still missing. Dislocations, whether

determined by substrate lattice mismatch, coalescing islands, or nanoparticles, seem

to play a determinant role in pinning vortices across the H, T phase space. However,

their direct observation, at least in the concentrations required to explain the high J

values measured in ﬁlms is still lacking. Moreover, the source of additional pinning

mechanisms such as the J

enhancement resulting from light rare earth substitu-

tion for yttrium is still not clear. Further developments in the understanding of the

physics and the phenomenology of high-T

superconductors are therefore necessary

for the ultimate success of coated conductors.

Acknowledgments The research presented in this chapter was sponsored by the US Department

of Energy, Ofﬁce of Electricity Delivery and Energy Reliability - Superconductivity Program, un-

der contract DE-AC05–00OR22725 with UT-Battelle, LLC managing contractor for Oak Ridge

National Laboratory.

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