An insight into what superconducts in polycrystallineboron-doped diamonds based on investigationsof microstructureN. Dubrovinskaia†‡§, R. Wirth¶, J. Wosnitza�, T. Papageorgiou�, H. F. Braun††, N. Miyajima‡‡, and L. Dubrovinsky††

†Mineralphysik und Strukturforschung, Mineralogisches Institut, Universitat Heidelberg, D-69120 Heidelberg, Germany; ‡Lehrstuhl fur Kristallographie,††Physikalisches Institut, and ‡‡Bayerisches Geoinstitut, Universitat Bayreuth, D-95440 Bayreuth, Germany; ¶GeoForschungsZentrum Potsdam,Experimental Geochemistry and Mineral Physics, 14473 Potsdam, Germany; and �Hochfeld-Magnetlabor Dresden, ForschungszentrumDresden-Rossendorf, D-01314 Dresden, Germany

Edited by Ho-kwang Mao, Carnegie Institution of Washington, Washington, DC, and approved June 22, 2008 (received for review February 20, 2008)

The discovery of superconductivity in polycrystalline boron-dopeddiamond (BDD) synthesized under high pressure and high temper-atures [Ekimov, et al. (2004) Nature 428:542–545] has raised anumber of questions on the origin of the superconducting state. Itwas suggested that the heavy boron doping of diamond eventu-ally leads to superconductivity. To justify such statements moredetailed information on the microstructure of the composite ma-terials and on the exact boron content in the diamond grains isneeded. For that we used high-resolution transmission electronmicroscopy and electron energy loss spectroscopy. For the studiedsuperconducting BDD samples synthesized at high pressures andhigh temperatures the diamond grain sizes are �1–2 �m with aboron content between 0.2 (2) and 0.5 (1) at %. The grains areseparated by 10- to 20-nm-thick layers and triangular-shapedpockets of predominantly (at least 95 at %) amorphous boron.These results render superconductivity caused by the heavy borondoping in diamond highly unlikely.

superconductivity � transmission electron microscopy

A fter the discovery of superconductivity in boron-dopeddiamond (BDD) (1), numerous theoretical and experimen-

tal studies (2–10) confirmed the phenomenon and went alongwith its explanation. Superconductivity in group IV semicon-ductors with diamond structure—such as silicon, germanium,and their alloys—was already predicted in the early 1960s tooccur at very low temperatures (11). Except for diamond (1–3,8, 10), there was a report on experimentally observed supercon-ductivity in boron-doped silicon (9). For diamond, Ekimov et al.(1) suggested that the superconducting transition temperature(Tc) increases with heavy boron doping (Tc � 4 K at 2.6 at % B).However, further investigations of superconducting BDD pre-pared by the high-pressure–high-temperature (HPHT) tech-nique and by chemical vapor deposition (CVD) revealed stronginhomogeneities in these materials (1, 3, 7, 10, 12). In particular,B-rich phases [such as B4C (1, 7, 10, 12), used as a reactant, andB50C2 (10, 12)] were found in HPHT BDD samples. In BDDprepared by CVD sp2-bonded carbon is unavoidable (3), and onecannot exclude that sp2-bonded amorphous or graphite-likecarbon accumulates some amount of boron. The presence ofextra phases and large discrepancies in the B content determinedby various methods (1–3, 7–10, 13, 14) (such as secondaryion-mass spectroscopy, Hall-effect measurements, correlationswith diamond lattice parameters, infrared spectroscopy, andmicroprobe analysis) and the absence of a clear correlationbetween Tc and the B concentration (10) raise the question ofhow much boron indeed is incorporated into the diamondstructure of superconducting samples. It is remarkable that,although measurements of the temperature dependence of theresistance have been conducted on BDD single crystals withboron content ranging from �1019 to 2.7 � 1021 cm�3 (i.e., upto 1.53 at % B), no evidence of superconductivity was found

down to 0.5 K (15, 16). This is at odds with the results for thepolycrystalline materials and contrasts especially with one report(3), where superconducting onset temperatures as high as 7.4 Kwere observed for CVD BDD samples with even smaller boronconcentrations [9.4 � 1020 cm�3 (0.53 at %)] (3).

We investigated polycrystalline BDD samples that were syn-thesised at 20 (1) GPa and 2,300 (50) K (four samples) as wellas at 9.0 (5) GPa and 2,500 (50) K (one sample) in a 5,000-tonnepress by using a HPHT technique as described in refs. 10 and 12.A mixture of graphite (referred further to as 12C) or isotopicallypure amorphous carbon 13C and B4C in a ratio C:B � 13:1 (�7at % B) were used as starting materials. Synchrotron andin-house x-ray diffraction investigations revealed for all fivesamples studied that their main crystalline component (�99%)is diamond with a small amount of boron carbide B50C2. In somesamples residuals of the starting material B4C were found (10,12). Microprobe analysis of the samples revealed a boronconcentration of 2.6 at % (4.6 � 1021 cm�3) in the samples,whereas Hall-coefficient measurements gave a charge-carrierconcentration of 1.4 � 1021 cm�3, that is, three times as low asthe apparent B concentration in the material (supporting infor-mation (SI) Fig. S1). Both microprobe analysis and energy-dispersive x-ray (EDX) spectra did not show the presence ofother elements than carbon and boron in BDD samples.

The superconducting transition temperatures were deter-mined by means of electrical transport and specific heat mea-surements (Fig. S2 and Fig. S3), and were found to be in goodagreement with earlier reports (1, 7, 10).

The boron content in our samples was characterized bymeasurements of diamond lattice parameters, Raman spectros-copy, Hall-effect measurements, scanning electron microscopy,and microprobe analysis (10, 12). All of the methods mentionedabove give only sample-averaged results. In particular, no infor-mation on the spatial distribution of boron on the submicronlevel in the diamond samples was available. To investigate themicrostructure and the exact boron location in our supercon-ducting BDDs we used transmission electron microscopy(TEM). Electron transparent foils of BDD samples were pre-pared by means of focused ion beam (FIB) techniques (for

Author contributions: N.D., H.F.B., and L.D. designed research; N.D., R.W., J.W., T.P., H.F.B.,N.M., and L.D. performed research; N.D., R.W., J.W., T.P., and N.M. contributed newreagents/analytic tools; N.D., R.W., J.W., H.F.B., N.M., and L.D. analyzed data; and N.D. andL.D. wrote the paper.

The authors declare no conflict of interest.

This article is a PNAS Direct Submission.

§To whom correspondence should be addressed at: Mineralphysik und Strukturforschung,Mineralogisches Institut, University of Heidelberg, Im Neuenheimer Feld 236, D-69120Heidelberg, Germany. E-mail: natalia.dubrovinskaia@min.uni-heidelberg.de.

This article contains supporting information online at www.pnas.org/cgi/content/full/0801520105/DCSupplemental.

© 2008 by The National Academy of Sciences of the USA

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details, see Materials and Methods). Fig. 1 shows two bright-fieldTEM images of our polycrystalline sample, which consists ofmicrometer-size carbon grains with predominantly diamondstructure as revealed by x-ray and confirmed by electron-diffraction data. A number of grains from different samples wereinvestigated by using high-resolution TEM (HRTEM) (SI Text,Fig. S4). All investigated grains are separated by layers ofamorphous material along straight boundaries with thicknessesof 10–20 nm. The material constituting the layers along the grainboundaries also fills triangular-shaped pockets at the grainjunctions (Fig. 2A). With side lengths reaching 0.5 �m some ofthe triangular pockets have sizes comparable to some diamondgrains. The complete absence of any TEM diffraction contrast(uniform light-gray contrast during sample tilting) indicates theamorphous state of the filler material (Fig. 2 A).

To probe the chemical nature of the inter- and intragranularmaterial in our samples, we measured the electron energy-lossspectra (EELS) at boron K and carbon K ionization edges fornumerous pockets and grains (Fig. 2B). The EELS spectra (Fig.2B Lower) revealed that the pocket material contains �95%boron with a small amount of carbon (�5–6%, varying for thedifferent pockets). The appearance of some amorphous boronmaterial under high-pressure, high-temperature syntheses in theB–C–N–O system was reported in refs. 17 and 18. The authors(17, 18) provided a detailed investigation of boron K edges fromseveral �-rh B-bearing materials. They demonstrated that the BK edge energy-loss near-edge structure (ELNES) of �-rh B andthe products with the �-rh B structure exhibit an edge rich infeatures (17). The ELNES for the B K edge from �-rh B andstructurally related materials was divided into the �* (�195 eV)

Fig. 1. Bright-field TEM images of a polycrystalline BDD sample. (A) A micrometer-sized grain is separated from other grains by a clearly visible straightboron-rich boundary. (B) Twin boundaries and dislocations observed within a grain.

Fig. 2. The results of TEM and EELS investigations. (A) Bright-field TEM image of a BDD sample. The amorphous material constituting the layers along the grainboundaries also fills triangular-shaped pockets at the grain junctions. (B) Electron energy-loss spectra (EELS) at the boron K and carbon K ionization edges fromthe pockets (Lower) and the diamond grains (Upper) reveal that the pocket material consists of �95 at % boron with a small amount (�5–6 at %, varying indifferent pockets) of carbon. The amount of boron in the diamond grains does not exceed 0.5%. The �* peak due to sp2-bonded carbon is very common for thepure diamond investigated by TEM on carbon grid. A smaller beam size was selected to measure the intergranular pocket, thus resulting in lower intensity anda more ‘‘noisy’’ spectrum (Lower).

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and �* (195–210 eV) regions. The B K edge shapes from the �-rhB-bearing materials could be divided into three types (17): �-rhB, B6O, and B4C. However, the EELS spectra we obtained fromthe carbon-doped amorphous boron phase (with sharp featuresat �189 eV and 199 eV in the �* and �* regions, respectively;Fig. 2B Lower) at the B K edge are different from that known forcrystalline boron, boron carbides, nitrides, oxides, and boroncarbooxinitrides (17–20). This suggests that the intergranularboron does not form B12 icosahedra typical for the structure ofpure boron and boron carbides. The observed EELS spectra arealso different from the spectrum of amorphous boron carbide(21). An example for an elemental mapping using EELS is shownin Fig. 3. Practically all boron is concentrated in the grainboundaries and pockets.

The accurate investigation of the interior of the diamond-structured grains (Fig. 3) shows that boron is quite homogeneouslydispersed in a very small amount within the carbon grains. Boronmapping did not show any directional dependence in B concentra-tion within the diamond grains. To quantify the amount of boronin the well crystallized diamond part of the grains, we measured anumber of EELS spectra at various points (Fig. 2B Upper) Ourresults show consistently that the true amount of boron incorpo-rated into the diamond structure is between 0.2 (2) and 0.5 (1) at% (see details of quantitative EELS analysis in Materials andMethods). That is considerably less than the 0.8 at % estimated fromthe number of charge carriers determined by the Hall-effectmeasurements, and almost one order of magnitude less than theoverall amount of B (2.6 at %) as determined by electron micro-probe analysis (10, 12). This result is consistent with our observationthat boron is mainly localized in intergranular noncrystalline layersand pockets.

Our EELS data show as well that there are a few small (5–50nm) isolated spots with very high B concentration within thediamond grains. These spots have platelet-like shapes (Fig. S5).HRTEM revealed their crystalline structure and the electrondiffraction patterns showed reflections at 4.36 Å and 3.92 Åcharacteristic for B50C2, in agreement with synchrotron diffrac-tion data obtained for this sample (12).

Our observations render it highly unlikely that the observedsuperconductivity in BDD synthesized at high pressures and hightemperatures is actually related to the boron incorporated in thediamond. If so, then even �0.5 at % B would be sufficient toinduce superconductivity in diamond with onset temperatures of�2.4 K and transition widths of �1.4 K. Our heat capacitymeasurements confirm the bulk nature of the superconductivity.However, the total amount of material converted to the super-conducting state in our samples is of the order of 20% or less,consistent with other reports (7) (Fig. S3). If boron doping at alevel of 0.5 at % would actually be enough to generate super-conductivity in diamond, then, based on our observation of ahomogeneous doping of all diamond grains, one would expectthe majority of the sample to become superconducting (not only�20%). Moreover, as mentioned, presently available data forsingle-crystalline BDD contradict the results obtained for poly-crystalline samples. Indeed, superconductivity has not beenobserved resistively in single-crystalline diamonds with B con-tents up to 1.53 at % (15, 16). However, the diamond grains inour superconducting HPHT BDD samples are completely iso-lated from each other by a boron-rich amorphous phase. Con-sequently, this phase controls the superconducting transition atleast in the electrical transport data.

Besides the investigation of BDD produced by the HPHTtechnique we were able to study the microstructure of a super-conducting (Tc � 5.2 K) boron-doped thin film deposited on asingle-crystalline diamond (111) surface by use of the microwaveplasma-assisted chemical vapor deposition (MPCVD) method(22, 23). The sample was kindly supplied by Dr. Y. Takano(National Institute for Materials Science, Tsukuba, Japan). OurHRTEM analysis surprisingly revealed, instead of the expecteddiamond morphology, a graphite-like structure (24) with ahomogenous B distribution within the film according to EELSmapping (Fig. S6).§§ Consequently, also for this sample, noevidence for superconductivity in BDD was found.

Based on the results of our research, we can suggest a fewhypotheses for the explanation of the nature of superconductiv-ity in the investigated samples: (i) Undistorted diamond becomessuperconducting at B concentrations much lower than reported,but that strongly contradicts the results on single-crystal BDDsamples (15, 16). This also does not agree with the experimentalheat capacity data showing that only a small fraction of thesample becomes superconducting, although diamond forms itsmajor part. (ii) The graphite-like regions with comparativelymoderate B concentrations inside the grains become supercon-ducting. However, these innermost parts of diamond grains arenot connected to each other; this hypothesis cannot explainsuperconductivity observed in resistivity measurements. (iii) Theintergranular boron-rich material becomes superconducting. Inour opinion, the third hypothesis is most probable, becausecarbon-doped amorphous boron phase, which is filling all of theintergranular space in the polycrystalline samples, forms acontinuous net through the whole sample and its amount agrees

§§Besides our HRTEM investigations, we observed graphite reflections at 3.34 Å (002), 2.04Å (101), and 1.79 Å (102) in diffraction patterns obtained by use of high-brilliancein-house and synchrotron x-ray diffraction (XRD). Hoesch et al. (24) as well studied BDDthin films and observed a (002) reflection (d � 1.79 Å) forbidden for diamond. This hintsat the presence of graphite in their samples because the (102) graphite reflection appearsat d � 1.79 Å.

Fig. 3. Elemental B map obtained by using EELS and consisting of fourimages of an intergranular area. It shows four diamond grains, five grainboundaries, and two pockets. The vast majority of boron is concentrated in thegrain boundaries and pockets. Boron is quite homogeneously dispersed in avery small amount within the diamond grains.

Dubrovinskaia et al. PNAS � August 19, 2008 � vol. 105 � no. 33 � 11621

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with the estimations of a superconducting part of the samplebased on the heat capacity experimental data.

To prove that superconductivity in diamonds is caused by acertain amount of boron doping one would need to measure wellcharacterized single-crystalline BDD. Otherwise, the presenceof other complex phases in the available composite materialshampers any definite conclusion.

Materials and MethodsDetails of the samples synthesis and preparation are described in refs. 10and 12.

Transmission Electron Microscopy (TEM). Electron transparent foils were pre-pared by focused ion beam (FIB) techniques. This allows the preparation ofsite-specific TEM foils with typical dimensions of 15–20 �m width, �10 �mlength, and �0.150 �m thickness. Focused Ga ions have been used to sputtermaterial off the sample. A special device, called a selected carbon mill (SCM)significantly enhanced the sputtering process of diamond. In this processOH-containing MgSO4 was used, which was heated inside the SCM device.Water vapor was brought close to the ion beam by inserting a needle close tothe location where the foil was cut. The H2O decomposed by the ion beam andoxygen oxidized the diamond. Details of the FIB technique and the use of aSCM are described elsewhere (25).

TEM investigations were performed with a TECNAI F20 XTWIN transmissionelectron microscope operating at 200 kV with a field emission gun electronsource. The TEM is equipped with a Gatan Tridiem filter, an EDAX Genesis x-rayanalyzer with an ultrathin window and a Fishione high-angle annular dark-field detector (HAADF). The Tridiem filter was used for the acquisition of

energy-filtered images applying a 20-eV window to the zero-loss peak. EELSspectra of the different K edges (B, C) were acquired in the diffraction modewith a dispersion of 0.1 eV per channel and an entrance aperture of 2 mm. Theresolution of the filter was 0.9 eV at half-width of the full maximum of thezero-loss peak. The acquisition time was 1 s. Processing of the spectra (back-ground subtraction, removal of multiscattering, and quantification) was per-formed by using the DigitalMicrograph software package.

Details of Quantitative EELS Analysis. Removal of plural scattering by Fourier-ratio deconvolution by using the DM software package; t � lambda was in therange of 0.7 to 0.8 (thickness � electron mean free path); acceptance semi-angle used was 1.8 mrad; EELS quantification occurred by using the DMsoftware package (EELS Quantification); cross-sections were calculated byusing the Hartree–Slater model; background model was Power Law; forbackground and signal windows, we used default values of the software.

From our long-standing experience with the FIB technique we can excludethat the observed noncrystalline layers and pockets are produced during theFIB milling. Irradiation damage as a cause of the noncrystallinity can beeliminated as well, because after insertion of the sample into the TEM electronbeam focusing was avoided until the tilt experiments had been carried out.

EDX spectra were obtained in the scanning transmission mode (STEM) byusing the TIA software package of the TEM. Significant mass loss duringanalysis was avoided by scanning the beam in a preselected window (20 � 20nm or larger). The spot size was �0.1 nm, and the acquisition time was 60 s atan average count rate of 60–80 cps. This resulted in a counting error of about4–5% at a 3� level.

ACKNOWLEDGMENTS. We thank G. Eska and P. McMillan for useful discus-sions. Work at Bayreuth was supported by Deutsche Forschungsgemeinschaft(DFG) through DFG Priority Program 1236.

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