Saimm 201710 oct

Page 1

VOLUME 117

NO. 10

OCTOBER 2017

advanced metals initiative

17–19 October 2017

ourfuture future through through science our science

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The Southern African Institute of Mining and Metallurgy ) ) ) ) ) ) ! ! !" " Mxolis Donald Mbuyisa Mgojo President, Chamber of Mines of South Africa ! ! " !" " Mosebenzi Joseph Zwane Minister of Mineral Resources, South Africa Rob Davies Minister of Trade and Industry, South Africa Naledi Pandor Minister of Science and Technology, South Africa !" " S. Ndlovu !" " " A.S. Macfarlane " ! " !" " M.I. Mthenjane ! " !" " Z. Botha " " !" " C. Musingwini ! ! !" !"! J.L. Porter ! ! " "! V.G. Duke I.J. Geldenhuys M.F. Handley W.C. Joughin E. Matinde M. Motuku D.D. Munro

G. Njowa S.M Rupprecht A.G. Smith M.H. Solomon D. Tudor D.J. van Niekerk A.T. van Zyl

!" " "! N.A. Barcza R.D. Beck J.R. Dixon M. Dworzanowski H.E. James R.T. Jones G.V.R. Landman

R.G.B. Pickering S.J. Ramokgopa M.H. Rogers D.A.J. Ross-Watt G.L. Smith W.H. van Niekerk R.P.H. Willis

G.R. Lane–TPC Mining Chairperson Z. Botha–TPC Metallurgy Chairperson A.S. Nhleko–YPC Chairperson K.M. Letsoalo–YPC Vice-Chairperson ! ! "! Botswana DRC Johannesburg Namibia Northern Cape Pretoria Western Cape Zambia Zimbabwe Zululand

L.E. Dimbungu S. Maleba J.A. Luckmann N.M. Namate W.J. Mans R.J. Mostert R.D. Beck D. Muma S. Matutu C.W. Mienie

!!" " "! Australia: I.J. Corrans, R.J. Dippenaar, A. Croll, C. Workman-Davies Austria: H. Wagner Botswana: S.D. Williams United Kingdom: J.J.L. Cilliers, N.A. Barcza USA: J-M.M. Rendu, P.C. Pistorius

) *Deceased * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * *

W. Bettel (1894–1895) A.F. Crosse (1895–1896) W.R. Feldtmann (1896–1897) C. Butters (1897–1898) J. Loevy (1898–1899) J.R. Williams (1899–1903) S.H. Pearce (1903–1904) W.A. Caldecott (1904–1905) W. Cullen (1905–1906) E.H. Johnson (1906–1907) J. Yates (1907–1908) R.G. Bevington (1908–1909) A. McA. Johnston (1909–1910) J. Moir (1910–1911) C.B. Saner (1911–1912) W.R. Dowling (1912–1913) A. Richardson (1913–1914) G.H. Stanley (1914–1915) J.E. Thomas (1915–1916) J.A. Wilkinson (1916–1917) G. Hildick-Smith (1917–1918) H.S. Meyer (1918–1919) J. Gray (1919–1920) J. Chilton (1920–1921) F. Wartenweiler (1921–1922) G.A. Watermeyer (1922–1923) F.W. Watson (1923–1924) C.J. Gray (1924–1925) H.A. White (1925–1926) H.R. Adam (1926–1927) Sir Robert Kotze (1927–1928) J.A. Woodburn (1928–1929) H. Pirow (1929–1930) J. Henderson (1930–1931) A. King (1931–1932) V. Nimmo-Dewar (1932–1933) P.N. Lategan (1933–1934) E.C. Ranson (1934–1935) R.A. Flugge-De-Smidt (1935–1936) T.K. Prentice (1936–1937) R.S.G. Stokes (1937–1938) P.E. Hall (1938–1939) E.H.A. Joseph (1939–1940) J.H. Dobson (1940–1941) Theo Meyer (1941–1942) John V. Muller (1942–1943) C. Biccard Jeppe (1943–1944) P.J. Louis Bok (1944–1945) J.T. McIntyre (1945–1946) M. Falcon (1946–1947) A. Clemens (1947–1948) F.G. Hill (1948–1949) O.A.E. Jackson (1949–1950) W.E. Gooday (1950–1951) C.J. Irving (1951–1952) D.D. Stitt (1952–1953) M.C.G. Meyer (1953–1954) L.A. Bushell (1954–1955) H. Britten (1955–1956) Wm. Bleloch (1956–1957)

* * * * * * * * * * * * * * * * * * * * * * * *

*

*

* * *

*

*

H. Simon (1957–1958) M. Barcza (1958–1959) R.J. Adamson (1959–1960) W.S. Findlay (1960–1961) D.G. Maxwell (1961–1962) J. de V. Lambrechts (1962–1963) J.F. Reid (1963–1964) D.M. Jamieson (1964–1965) H.E. Cross (1965–1966) D. Gordon Jones (1966–1967) P. Lambooy (1967–1968) R.C.J. Goode (1968–1969) J.K.E. Douglas (1969–1970) V.C. Robinson (1970–1971) D.D. Howat (1971–1972) J.P. Hugo (1972–1973) P.W.J. van Rensburg (1973–1974) R.P. Plewman (1974–1975) R.E. Robinson (1975–1976) M.D.G. Salamon (1976–1977) P.A. Von Wielligh (1977–1978) M.G. Atmore (1978–1979) D.A. Viljoen (1979–1980) P.R. Jochens (1980–1981) G.Y. Nisbet (1981–1982) A.N. Brown (1982–1983) R.P. King (1983–1984) J.D. Austin (1984–1985) H.E. James (1985–1986) H. Wagner (1986–1987) B.C. Alberts (1987–1988) C.E. Fivaz (1988–1989) O.K.H. Steffen (1989–1990) H.G. Mosenthal (1990–1991) R.D. Beck (1991–1992) J.P. Hoffman (1992–1993) H. Scott-Russell (1993–1994) J.A. Cruise (1994–1995) D.A.J. Ross-Watt (1995–1996) N.A. Barcza (1996–1997) R.P. Mohring (1997–1998) J.R. Dixon (1998–1999) M.H. Rogers (1999–2000) L.A. Cramer (2000–2001) A.A.B. Douglas (2001–2002) S.J. Ramokgopa (2002-2003) T.R. Stacey (2003–2004) F.M.G. Egerton (2004–2005) W.H. van Niekerk (2005–2006) R.P.H. Willis (2006–2007) R.G.B. Pickering (2007–2008) A.M. Garbers-Craig (2008–2009) J.C. Ngoma (2009–2010) G.V.R. Landman (2010–2011) J.N. van der Merwe (2011–2012) G.L. Smith (2012–2013) M. Dworzanowski (2013–2014) J.L. Porter (2014–2015) R.T. Jones (2015–2016) C. Musingwini (2016–2017)

! !($( ) ' $ ) # &"'(" Scop Incorporated #&%!(" Genesis Chartered Accountants ' ('%$(&'" The Southern African Institute of Mining and Metallurgy Fifth Floor, Chamber of Mines Building 5 Hollard Street, Johannesburg 2001 • P.O. Box 61127, Marshalltown 2107 Telephone (011) 834-1273/7 • Fax (011) 838-5923 or (011) 833-8156 E-mail: journal@saimm.co.za

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VOLUME 117

#&%!(&$ ) !$(# R.D. Beck P. den Hoed M. Dworzanowski B. Genc M.F. Handley R.T. Jones W.C. Joughin J.A. Luckmann C. Musingwini S. Ndlovu J.H. Potgieter T.R. Stacey M. Tlala

VOLUME 117 NO. 10 OCTOBER 2017

Contents Journal Comment: by G. Pattrick . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

iv

President’s Corner: by S. Ndlovu . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Conference Paper Number—AMIPM–S05 Reaction kinetics of ZrF4 chloridation at elevated temperatures by N.J.M. Grobler, C.J. Postma and P.L. Crouse . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

927

# '(%&"& ) ' ('"' %$%& '

Conference Paper Number—AMIPM–S08 Design and fabrication of a chemical vapour deposition system with special reference to ZrC layer growth characteristics by S. Biira, P.L. Crouse, H. Bissett, T.T. Hlatshwayo, J.H. van Laar and J.B. Malherbe . . . . . . . . . . .

931

Barbara Spence Avenue Advertising Telephone (011) 463-7940 E-mail: barbara@avenue.co.za

Conference Paper Number—AMIPM–S11 Selective sublimation/desublimation separation of ZrF4 and HfF4 by C.J. Postma, S.J. Lubbe and P.L. Crouse . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

939

The Secretariat The Southern African Institute of Mining and Metallurgy

Conference Paper Number—AMIPM–S13 Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb by S.S. Mahlalela and P.G.H. Pistorius . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 947

#&%!(&$ ) ! " %$ % D. Tudor

'"'%)$ #) &" '#) The Southern African Institute of Mining and Metallurgy P.O. Box 61127 Marshalltown 2107 Telephone (011) 834-1273/7 Fax (011) 838-5923 E-mail: journal@saimm.co.za

(& %'#) ) Camera Press, Johannesburg

ISSN 2225-6253 (print) ISSN 2411-9717 (online)

Conference Paper Number—AMIPM–S23 Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst for aquacatalytic Sonogashira cross-coupling reaction by P.P. Mpungose, N.I. Sehloko, T. Cwele, G.E.M. Maguire and H.B. Friedrich. . . . . . . . . . . . . . . . .

955

THE INSTITUTE, AS A BODY, IS NOT RESPONSIBLE FOR THE STATEMENTS AND OPINIONS ADVANCED IN ANY OF ITS PUBLICATIONS.

Conference Paper Number—AMIPM–S08 First-principles studies of Fe-Al-X (X=Pt, Ru) alloys by C.S. Mkhonto, H.R. Chauke and P.E. Ngoepe . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

963

CopyrightŠ 2017 by The Southern African Institute of Mining and Metallurgy. All rights reserved. Multiple copying of the contents of this publication or parts thereof without permission is in breach of copyright, but permission is hereby given for the copying of titles and abstracts of papers and names of authors. Permission to copy illustrations and short extracts from the text of individual contributions is usually given upon written application to the Institute, provided that the source (and where appropriate, the copyright) is acknowledged. Apart from any fair dealing for the purposes of review or criticism under The Copyright Act no. 98, 1978, Section 12, of the Republic of South Africa, a single copy of an article may be supplied by a library for the purposes of research or private study. No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means without the prior permission of the publishers. Multiple copying of the contents of the publication without permission is always illegal.

Conference Paper Number—KN02 PGMs: A cornucopia of possible applications by L.A. Cornish. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

969

Conference Paper Number—AMIPM–MS35 Metal and alloy spheroidisation for the Advanced Metals Initiative of South Africa, using high-temperature radio-frequency plasmas by H. Bissett and I.J. van der Walt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

975

Conference Paper Number—AMIPM–MS39 A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys in the as-cast and annealed conditions by B.O. Odera and L.A. Cornish . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

981

Conference Paper Number—AMIPM–MS44 Alternative carbon materials as practical and more durable fuel cell electrocatalyst supports than conventional carbon blacks by M.L. Stevenson and G. Pattrick . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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%'( $%&! $ ) # &"!( ) !$(# VOLUME 117

R. Dimitrakopoulos, McGill University, Canada D. Dreisinger, University of British Columbia, Canada E. Esterhuizen, NIOSH Research Organization, USA H. Mitri, McGill University, Canada M.J. Nicol, Murdoch University, Australia E. Topal, Curtin University, Australia

VOLUME 117

N O. 10

OCTOBER 2017

advanced metals initiative

All papers featured in this edition were presented at the

The Precious Metals Development Network Conference (PMDN) 17–19 October 2017

ourfuture future through through science our science

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t is with great pleasure that we once again present the annual conference of the Advanced Metals Initiative (AMI). The AMI was established jointly by the Department of Science and Technology (DST) and the science councils, namely, Mintek, NECSA and the CSIR and has received generous funding from the DST since 2003. The principal objective of the AMI is to increase local industrialization through a research and development led initiative. In essence the AMI focuses on the downstream beneficiation of local resources and aims to achieve this through the development of materials, applications and technologies that enable the formation of new industries, enhance the competiveness of existing industries or localise the production of existing advanced and critical products. The AMI is split into four networks, each focused on specific local resources. We have the Precious Metals Development Network (PMDN) active in platinum group metals and gold, the Light Metals Development Network (LMDN) looking at aluminium and titanium, the Ferrous Metals Development Network (FMDN) considering iron and alloying base metals, and the Nuclear Metals Development Network (NMDN) advancing zirconium, hafnium and tantalum processes and products. A critical aspect of an initiative like the AMI is the development of the highly skilled human capital ultimately required to implement, maintain and drive the envisioned industries that could result from the programme. Here the AMI has made a significant contribution by the support and graduation of over 140 postgraduates (PhD and MSc levels) at universities across South Africa in projects and programmes that are specifically aligned with the aims of each network. Each year several of the currently enrolled students are afforded the opportunity to present their work in the specifically set aside Student Day of the annual conference. Several of these excellent papers will also be included in the issue of the SAIMM Journal issue dedicated to the conference. It is Mintek’s honour this year to host the AMI Precious Metals 2017 conference, covering topics specific to the PMDN in the areas of chemistry, catalysts and materials as they relate to potential downstream value addition of our local precious metals. An exciting development this year is that we have partnered with the SAIMM Platinum 2017 conference. Thus in one venue and in parallel sessions we have an event that will cover the entire value chain of platinum group metals from new developments in exploration and mining to advanced materials and applications of the metals. It is hoped that delegates at the conference will use the occasion to explore the full event and gain insight into the extensive activities undertaken in South Africa and recognise the opportunity for industrial growth. The Organising Committee takes this opportunity to welcome delegates to the conference and to thank authors and reviewers for their contributions. We also thank Camielah Jardine and her team at SAIMM for their excellent organisation of the conference. In addition we thank Llanley Simpson and Wilna du Plessis from the DST for their guidance and for generous funding of this conference and the AMI programme.

G. Pattrick Chairperson of the Organizing Committee

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would like to welcome you all to the new 2017/2018 term of the SAIMM. I hope you are all as excited as I am about what lies ahead. I am greatly looking forward to spending the next twelve months with you as we discuss and delve into different topics that interest us as members of this marvellous Institute. As you are all aware, the SAIMM was established in 1894 and will be turning 125 years of age in 2019. The Institute has been built on a strong foundation of selfless commitment and diligence from its members; mostly on the basis of volunteerism. I would therefore, like to start off the year by acknowledging and thanking all the members of the SAIMM who have selflessly given of themselves and their time in the past, and will hopefully continue to do so in the future to keep the mission and goals of the Institute alive and active. It is indeed not easy, especially as we try to balance our demanding careers and personal lives. However, it does add value to our worth as we contribute to a bigger goal than our individual selves. The SAIMM is built on the foundations of the mining and minerals industry. The Institute serves the interests of the professionals that work in the industry and has grown from strength and strength due to the support of the mining industry and related companies from which most of our members are drawn. It is interesting to note that although the mining and minerals landscape has changed from what it was 120 years ago, we still continue to receive a lot of support from the mining industry. We should not take this support lightly, but be more cognizant of it and the value that it embodies. It also interesting to note that the focus of the SAIMM, which has always been to keep the interests of its members at the forefront, has not changed since the inception of the institute in 1894. However, this does not mean that the SAIMM has remained stagnant. If this was the case, it would not have survived to the present day this age. In my Presidential Address, I alluded to the fact that standing still is the same as going backwards and that companies that do not innovate get left behind and subsequently perish. Since we have made it to this age, it is evident that this cannot be said of our Institute. We have never stood still; we have continued to change while keeping our key mission and goals in focus. This has not always been easy. I am sure that if we delve into the history of the SAIMM and chat to those who have been part of the Institute over a long period, we would find that the Institute has faced numerous challenges in the past. Since the Institute is built on the foundations of the mining and minerals industry, such challenges should be expected. Just as the mining industry goes through cyclic downturns, it is natural for the SAIMM to go through challenges as well. And just as companies have had to reassess and re-evaluate their processes, structures, and operating procedures in order to remain competitive in the market, so the SAIMM has also had to make sure it had a clear and comprehensive grasp of external challenges as well as solutions in order to survive. Not only that, the SAIMM has also developed the capacity to see what others cannot see and thus turned such insight into opportunities. The Institute has had a comprehensive assessment of its strengths and limitations and has been committed to change where it has been relevant and appropriate. Where has all this come from you might ask? The source of the Institute’s strength has always been its members. Members who have shown an interest in the Institute’s activities have thus forced the Institute to always reassess and re-evaluate its offerings to its members as well as its standing as a professional organization. What is my key message in all this? It is that we should never sit back, but always strive for more. We don’t want to become extinct, and we should never be afraid of change; change should be seen as a chance for new opportunities. Lastly, we should always remember that we are an amazing organization; not many organizations make it to 123 years of age, and we have defied the ravages of time. More significantly, the only reason we are this old is because of actively participating members. So, as your President, I say thank you, dankie, ngiyabonga, enkosi, kea leboha, ndo livhuwa!

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S. Ndlovu President, SAIMM

DST PROFILE The vision of the Department of Science and Technology (DST) is to create a prosperous society that derives enduring and equitable benefits from science and technology. Its mission is to develop, coordinate and manage a national system of innovation that will bring about maximum human capital, sustainable economic growth and improved quality of life for all. When the DST was established in 2004, it was assigned a leading role in modernising South Africa's economy. The 2002 National Research and Development Strategy (NRDS) was well received in the science system, and substantial financial resources invested by the government, partners and stakeholders have seen significant progress in the attainment of the strategy's objectives. These include an integrated approach to human capital development, investment in science and technology infrastructure, and improving the strategic management of the public science and technology system to optimise its contribution to wealth creation and an improved quality of life. Another important DST strategy, the Ten-Year Innovation Plan (2008-2018) was developed to help drive South Africa towards sustained economic growth and a knowledge-based economy. The Plan introduces five "grand challenges" that continue to receive priority attention from government and researchers, to find ways of bridging the "innovation chasm" between research results and socio-economic benefits. For this to be achieved, government through the DST has been making important investments in human capital development, knowledge production, and the exploitation of such knowledge for commercial purposes. The five "grand challenges" include strengthening the bioeconomy, growing space science and technology, assuring energy security, responding to global change (with a focus on climate change), and contributing to a global understanding of shifting social dynamics in development. A quick look at the events and milestones of the past few years reveal that the country is positioning itself among global leaders in innovation, science and technology. One of the most iconic indicators in this regard is the country's remarkable work as

part of the global science project to build the world's largest radio telescope, the Square Kilometre Array, much of which is being hosted in the Northern Cape. The project has attracted considerable financial investment in infrastructure and skills advancement. In additive manufacturing, DST-funded work is driving and funding the national development and adoption of metal and polymer additive manufacturing technology to position South Africa as a competitor in the global market. Through the Collaborative Programme in Additive Manufacturing and the Aeroswift Technology Platform, it is developing and improving design capabilities specifically for the additive manufacturing industry, as well as building a platform for the production of large titanium structures, with a very specific focus on the aerospace industry. With science and technology critical elements in the growth and development of Africa, the DST is also playing a leading role in the development, integration and unification of the continent. The Department continues to strengthen its science, technology and innovation relations through a diverse portfolio of bilateral and multilateral relations, including with the BRICS grouping (Brazil, Russia, India, China and South Africa) and the European Union. South Africa is also building muscle in the areas of health research, earth observation, renewable energy and the beneficiation of metals. Information and communication technology are being used in pilot projects to show how technologies can benefit South Africans, especially in previously marginalised communities, in respect of education needs, among others. The DST's Technology Localisation Programme is being used to improve the competitiveness of small and medium enterprises. As required by the National Development Plan, the DST will continue its efforts to provide South Africa with enduring benefits from science, technology and innovation. For more information contact Zama Mthethwa, Account Executive. Zama.Mthethwa@dst.gov.za or www.dst.gov.za

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a1

Reaction kinetics of ZrF4 chloridation at elevated temperatures by N.J.M. Grobler*, C.J. Postma†and P.L. Crouse*

9 30)/1/ Zirconium metal is used in the alloy cladding of nuclear fuel rods. Necsa produces ZrF4 by reacting ammonium bifluoride with desilicated plasmadissociated zircon. The ZrF4 then undergoes a sublimation separation step to reduce the Hf content to below 100 ppm Hf. The ZrF4 is converted into ZrCl4 via its reaction with magnesium chloride. The ZrCl4 is then used in a plasma process to produce Zr metal for use in the nuclear industry. The aim of the research reported here is to obtain a first-order estimate of the reaction kinetics for the chloridation reaction. This was done using the data obtained from a dynamic thermogravimetric analysis experiment and minimising the error between experimentally determined degrees of conversion and predictions of a stepwise kinetic model predictions. The reaction was found to take place only above the sublimation point of ZrCl4. At a low enough heating rate it can be assumed the loss in mass is due to sublimation of ZrCl4, the moment it is formed and the rate of mass loss is equal to the reaction rate. 6 02./ reaction kinetics, zirconium chloride, thermogravimetric analysis.

83520.-*5103 Zirconium is used in the nuclear industry in the cladding of nuclear fuel rods. Zirconium naturally occurs with 1–3 % hafnium; nuclear grade zirconium needs to be purified to contain less than 100 ppm hafnium (Smolik, Jakbik-Kolon and Poraski, 2009). Part of the process developed by Necsa is to convert ZrF4 to ZrCl4 after the sublimation purification step (Postma, Niemand and Crouse, 2015). Makhofane et al. (2013) described a method for the chloridation of ZrF4 according to Reaction [i] below:

Ltd and were prepared through the method discussed by Nel et al. (2011). Each reagent was carefully weighed with a 10 % molar excess of MgCl2. The reagent mixture was roughly 6.4 mg ZrF4 and 8 mg MgCl2. Aluminum crucibles were used. The reagents were prepared under atmospheric conditions due to the non-availability of a glove box. The MgCl2 thus became partially hydrated due to moisture in the air. The degree of hydration was calculated from the first decomposition step where two H2O molecules are released. The TGA was programmed to heat at a constant rate of 1 °C per minute from room temperature to 500 °C and remain there for 10 minutes before cooling down. The analysis was done with 200 mL¡min-1 N2 as carrier gas. A Hitachi STA 7300 TGA-DTA was used for the work. The code for data fitting was developed inhouse and written in Python. It comprised prediction of the sample¡mass as a function of time and temperature by integration of the rate expressions using estimated kinetic parameter values, then minimising the difference between predicted and experimental values by refinement of the parameters.

:6/-(5/743.7.1/*-//1037 Figure 1 shows the dynamic thermogram of the reaction. MgCl2 is hygroscopic and hydrates readily under ambient conditions. Since we did not have the equipment to work under anhydrous conditions, we had to work with wet material

[i]

)621,6354( The materials were obtained from Necsa SOC

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* Fluoro-Materials Group, Department of Chemical Engineering, University of Pretoria, South Africa. †South African Nuclear Energy Corporation (Necsa), South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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The objective of this research is to estimate the reaction kinetics for the solid state chloridation of ZrF4 sufficient for scale-up and reactor design. The reagents were prepared with a slight molar excess of MgCl2 and dynamic thermogravimetric analysis (TGA) at a very low temperature ramp rate to ensure minimal thermal lag, was used to monitor the reaction.

Reaction kinetics of ZrF4 chloridation at elevated temperatures

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While the focus of this work is the chloridation reaction itself, the dehydration had of necessity to be incorporated to estimate the amount of dry MgCl2 present. The kinetic data given by Huang et al. (2010) was used for ease of implementation. Their kinetic models for the individual steps were adopted without change, while their pre-exponential factors and activation energies were used as starting values for obtaining a fit to our data. The first five dehydration steps, numbered 1–5 in Figure 1, correspond to Reactions [ii] to [vi] and to Equations [2] to [6] below. The chloridation of ZrF4, which starts at 396°C, is marked as step 6, corresponding to Reaction [i] and to Equation [1]. This is followed by a further loss of HCl as the hydroxide chloride, marked as step 7 and corresponding to Reaction [vii] and Equation [7]. The fitting procedure involved minimising the error between the model predictions and the experimental data, by adjusting the pre-exponential constant in the kinetic equations for the six decomposition steps and in the chloridation reaction itself, as well as the exponent in Equation [1].

and included this in our data analysis. Steps 1–5 and step 7 in Figure 1 are accounted for by dehydration, while step 6 is the actual chloridation reaction. An important assumption for the model is that the magnesium chloride starting material is a mixture of anhydrous MgCl2 and fully hydrated MgCl2¡6H2O and that the chloridation is effected by anhydrous MgCl2 rather than the hydrated material, or the oxychloride that forms during the dehydration of the hexahydrate. MgCl2¡6H2O loses crystal water upon heating through several decomposition steps (Kirsh, Yariv and Shoval, 1987; Huang et al., 2010). The dehydration of MgCl2¡6H2O is described by Reactions [ii] to [vii] below:

[1]

[2]

[3]

[4]

[5]

[ii] [iii]

[6]

[iv]

[7] In the above , the extent of reaction, is given by

[v]

[vi] [vii] It is important to note that MgO, not MgCl2, is the final product. The formation of gaseous product is reflected on the thermograms as a loss of mass. Kirsh, Yariv and Shoval (1987) and Huang et al. (2010) derived kinetic data for the thermal decomposition of MgCl2¡6H2O. It should be noted that neither was done according to most recent ICTAT specifications (Vyazovkin et al., 2011) and the kinetic triplets reported should be regarded as estimates rather than correct absolutes. In this work we similarly aim at obtaining engineering values, rather than exact kinetics, for the chloridation reaction.

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[8] where m0 and mf are the initial and final masses, is the heating rate, ki the pre-exponential factors and Eai the activation energies. The form of the dehydration equations given by Huang et al. (2010) can be attributed to the solid-state process mechanism used to fit the decomposition steps. The mechanisms that resulted in the best fit for each step are: two-dimensional phase boundary mechanism, threedimensional phase boundary mechanism, nucleation and nuclei growth mechanism (Avrami-Erofeev equation n=3), two-dimensional phase boundary mechanism, threedimensional diffusion mechanism (cylinder and G-B equation) and nucleation and nuclei growth mechanism (Avrami-Erofeev equation n=1), respectively. The general form of the kinetic data given by Huang et al. (2010) is shown in Equation [9]:

Reaction kinetics of ZrF4 chloridation at elevated temperatures [9] where f( ) is termed the kinetic model, for a single step and is given by the above-mentioned mechanisms for solid-state processes. For the chloridation reaction we assumed the simple empirical model, f( )= n. Table I shows the values of the pre-exponential constant, the activation energy and the temperature ranges during which each stage is active. The value of n, the exponent in the kinetic model for the chloridation step, Reaction [i], stage 6 was found to be 1.5. Figure 2 shows the experimental data along with the fitted kinetic model. The Euclidian norm of the difference between the experimental data and model prediction was 0.018, indicating a good approximation. In the model each was assigned a range 0–1. The initial and final masses of each were estimated from the thermogram. These mass values correspond to the temperature ranges for each step listed in Table I. The thermodynamically computed vapour pressure of zirconium tetrachloride is given in Figure 3. The inflection point of the sublimation curve is at 327°C. The onset temperature for the chloridation process, step 6, is 396°C. This suggests that the reaction is more complex than a typical solid-solid reaction where contact and mass transfer are problematic, causing very slow reaction rates. The chloridation takes place rapidly at a temperature where gaseous zirconium tetrachloride is the thermodynamically preferred phase and this reaction product can be released as soon as it is formed.

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!03*(-/103/743.726*0,,63.45103/ The work presented here yielded fair kinetic data, sufficient for reactor and process design. It further addresses the very real prospect that industrially it may be extremely difficult to use pure anhydrous MgCl2 as reagent, due to its hygroscopic nature. In order to make the research scientifically more rigorous, a number of issues need to be addressed in future work. The most important of these are: Experimental work under isothermal conditions are required for exact kinetic data Anhydrous kinetics has to be decoupled from the kinetics of the reaction of ZrF4 with the hydrated species, using pure compounds The exact hydrated magnesium species formed have to be determined as function of degree of bulk hydration

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Physical examination of the product at various stages during the reaction has to be done, by e.g. electron microscopy, to determine the mass transfer mechanisms that apply The kinetic triplets for the dehydration steps, reported in the literature and used here, need to be revisited. The pre-exponential factors and activation energies are unrealistically high in general and more extensive experimental work is required to achieve physically realistic values.

Table I

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Stage 1 Equation (2) Stage 2 Equation (3) Stage 3 Equation (4) Stage 4 Equation (5) Stage 5 Equation (6) Stage 6 Equation (1) Stage 7 Equation (7)

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Reaction kinetics of ZrF4 chloridation at elevated temperatures zirconia-based material with ammonium bi-fluoride. US patent 8778291 B2. South African Nuclear Energy Corporation Limited.

:6'6263*6/ HUANG, Q.Z., LU, G.M., WANG, J. and YU, J.G. 2010. Mechanism and kinetics of thermal decomposition of MgCl2.6H2O. Metallurgical and Materials Transactions B, vol. 41, no. 5. pp. 1059–1066. KIRSH, Y., YARIV, S. and SHOVAL, S. 1987. Kinetic analysis of thermal dehydration and hydrolysis of MgCl26H2O by DTA and TG. Journal of Thermal Analysis, vol. 32, no. 2. pp. 393–408. MAKHOFANE, M.M., NEL, J.T., HAVENGA, J.L. and AFOLABI, A.S. 2013. Chlorination of anhydrous zirconium tetrafluoride with magnesium chloride. Proceedings of the Precious Metal Development Network Conference, Cape Town, South Africa, 14–16 October 2013. Southern African Institute of Mining and Metallurgy, Johannesburg. pp. 339–348 NEL, J.T., DU PLESSIS, W., CROUSE, P.L. and RETIEF, W.L. 2011. Treatment of

POSTMA, C.J., NIEMAND, H.F. and CROUSE, P.L. 2015. A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form. Journal of the Southern African Institute of Mining and Metallurgy, vol. 115, no. 10. pp. 961–965. SMOLIK, M., JAKBIK-KOLON, A. and PORASKI, M. 2009. Separation of zirconium and hafnium using diphonix chelating ion-exchange resin. Hydrometallurgy, 95 3–4, pp. 350-353. VYAZOVKIN, S., BURNHAMB, A.K., CRIADOC, J.M., PÉREZ-MAQUEDAC, L.A., POPESCUD, C. and SBIRRAZZUOLIE, N. 2011. ICTAC Kinetics Committee recommendations for performing kinetic computations on thermal analysis data. Thermochimica Acta, vol. 520, nos. 1–2. pp. 1–19.

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http://dx.doi.org/10.17159/2411-9717/2017/v117n10a2

Design and fabrication of a chemical vapour deposition system with special reference to ZrC layer growth characteristics by S. Biira*†, P.L. Crouse*, H. Bissett‥, T.T. Hlatshwayo*, J.H. van Laar* and J.B. Malherbe*

The overall aim of this research project was to design and construct an inhouse, thermal chemical vapour deposition (CVD) reactor system, operating at atmospheric pressure. Radio frequency (RF) induction heating was used as the energy source, with a vertical-flow design, using thermally stable materials. The steps in the design and construction of this CVD system are described in detail. The growth conditions at different substrate temperatures, gas flow ratios and substrate-gas inlet gaps were assessed as part of the project. The growth rate of ZrC layers increases with increasing substrate temperature. The microstructure properties of the ZrC layers such as lattice parameters and orientation of crystal planes were all found to be dependent on deposition temperature. The increase in free carbon in the as-deposited coatings as the temperature increased was found to be a stumbling block for obtaining stoichiometric ZrC coatings. The surface morphology of the as-deposited ZrC layers also depends on the deposition parameters.

:& 57*3 ZrC, chemical vapour deposition, RF heating.

4975*029854 Chemical vapour deposition (CVD) systems are used in applications that require the deposition of layers or coatings of a material onto a surface (Park and Sudarshan, 2001). CVD plays an important role in the fabrication of microelectronic devices and in the formation of protective coatings (Pierson, 1999). CVD processes have been widely used commercially; however, the detailed design and optimisation of inexpensive CVD equipment and processes are not well described. Therefore the objectives of this study were; (1) to design and fabricate a CVD system, (2) to establish the operating conditions for achieving deposition in CVD mode and (3) then to deposit ZrC layers from ZrCl4 and methane. Zirconium carbide (ZrC) is very hard interstitial refractory carbide (Pierson, 1996) which has been proposed as a possible new material to be used as a protective coating layer for nuclear fuel particles because of its excellent properties (Meyer Fielding and Gan, 2007; Katoh et al., 2013). For example, ZrC has a very high melting temperature of about 3540°C and low density of 6.59 g/cm3 compared to other refractory carbides. It is chemically inert to various reagents and highly

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<08.+:49;/6%78269854 ;254397029854;64* 7(;*:.5389854 Prior to the designing of the CVD system for the deposition of ZrC layers shown in Figure 1, a thermodynamic analysis was carried out. This was done in order to establish the feasibility of the chemical reactions both at room conditions and at elevated temperatures. The thermodynamic analysis results were used as guidelines for the selection of the appropriate materials for the reactor system. The zirconium and carbon sources source materials were ZrCl4 and CH4 respectively. The feasibility of the chemical reaction relies on the change in Gibbs free energy. When the Gibbs free energy of the reaction is negative, then the reaction is feasible at the given conditions (Yan and Xu, 2010). The Gibbs free energy of formation and the corresponding equilibrium thermodynamic constant were calculated. The results are given in a previous publication (Biira et al., 2017b). HSC Chemistry software (a thermochemical software package) was also used to show the speciation of the reactants with temperature (Biira et al., 2017b). Based on the results of the thermodynamic analysis, deposition temperatures ranging from 1200°C to 1600°C were selected.

* University of Pretoria, South Africa. †Busitema University,Uganda.

‥ The South African Nuclear Energy Corporation (Necsa), Pretoria, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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&45.383

resistant to corrosion and wear (Pierson, 1996; Won et al., 2007). Many attempts have been made to grow ZrC layers; however, growing polycrystalline ZrC layers with good stoichiometry has remained a challenge. CVD has been identified as a suitable technique for the deposition of layers with relatively high purity and uniformity levels (Yan and Xu, 2010).

Design and fabrication of a chemical vapour deposition system

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In this study a vertical-wall thermal CVD system, as indicated in Figure 1, was developed in-house at the South African Nuclear Energy Corporation (Necsa) SOC Ltd. This inhouse-built CVD system consisted of four basic components. These were (i) the gas supply and delivery system (including ZrCl4 vaporiser system), (ii) an RF power supply system and induction coil, (iii) the reactor system (including the graphite reaction chamber) and (iv) exhaust/scrubber system.

workpiece). Table I gives the summary of the specifications of the coil used. To ensure maximum uniform heating, the effective heating length of the coil on the graphite reaction chamber was estimated using Equation [2] (Wai, Aung and Win, 2008) and was found to be 6.34 cm.

! ! ! !

where hg is the portion of the length of the cylindrical graphite tube with maximum heating, Cp is the pitch of the coil winding, dc is the diameter of conductor (copper tubing) and N is number of coil turns able to produce maximum uniform heating on about 6.34 cm of the length of the graphite reaction chamber. The substrate temperatures were measured by an infrared optical pyrometer through a quartz viewing window at the top of the flange (see Figure 2). The power calibration experiments were carried out five times and the average temperature values were considered to avoid errors that may arise from the pyrometer readings.

An RF induction heating source with maximum power output of 10 kW was used. This RF power supply model EASYHEAT LV AT 8310 was manufactured by Ambrell, Ameritherm Inc., USA. It was connected to a copper tubing induction heating coil. The coil was made of copper in order to minimise the heat losses and to maximise heating efficiency. The copper coil was water-cooled to prevent melting. The gap between the four-turn helical copper coils and the graphite reaction chamber was kept to a minimum to ensure good RF coupling. This distance was sufficient to accommodate the thermal insulation (using ceramic material) to reduce heat transfer from the heated graphite reaction chamber to the coil. The ceramic material also secured the graphite reaction chamber at the centre of the coil in a fixed position. When the material to be heated is placed close to the coil, the current increases, thus increasing the amount of heat induced in material. To achieve uniform heating the coupling distance Cd should be equal to, but not more than, twice the coil pitch (Simpson 1960), since heating effects also depends on the pitch of coil windings. The coupling distance was estimated using Equation ([1]): [1] where Din is the inner diameter of the coil and dg is the diameter of the graphite reaction chamber (acting as the

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[2]

Table I

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Material Coil type Number of turns Inner diameter (cm) Outer diameter (cm) Diameter of the coil tubing (cm) Coil length (cm) Coil pitch (cm) Coupling distance (cm) Gap between the windings

Copper Helical (round) 4 5.1 6.38 0.64 3.60 0.945 0.80 0.305

Design and fabrication of a chemical vapour deposition system

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! ! The CVD reactor system shown in Figure 1 was fabricated and constructed at Necsa. The elbow stainless steel outer casing and the graphite reaction chamber are among the many components. The schematic of the CVD reaction chamber system and components is given in Figure 2. One of the roles of the stainless steel outer casing was to confine the corrosive exit gases and enable safe extraction. The cylindrical high-density graphite tube manufactured by Le Carbone-Lorraine had an inner diameter of 2.5 cm and a length of 30 cm. The cylindrical graphite tube (reaction chamber), together with a substrate stage made of the same material attached at its bottom end, was fixed inside the steel casing (see Figure 2). The 74.1 g solid cylindrical graphite substrate stage had a length of 13.5 cm and diameter of 2.2 cm. This graphite tube was at the same time used as the workpiece for induction heating to attain the required temperature for the substrates (1200°C up to 1600°C for this experiment). Table II gives the summary of the material specifications and properties of the graphite tube that was used in this study. The graphite tube was equipped with leads on both ends, i.e. the gas inlet at the top and outlet at the bottom. The gas inlet and outlet permitted the flow of the gases past the heated substrates mounted on substrate stage. The distance from the inlet to the substrate stage varied between 70 mm and 170 mm. Zirconia wool was used for both thermal and electrical insulation and was placed between the steel casing and the graphite reaction chamber.

depressurising it in order to remove air (oxygen) and atmospheric moisture. Methane, hydrogen and argon flow rates were measured by calibrated rotameters and directed into the reaction chamber as shown in Figure 1 and Figure 3. The deposition was carried out at atmospheric pressure (which was about 87 kPa absolute at the location). The volumetric flow rates of argon, hydrogen and methane were controlled manually by pre-calibrated flow meters. Argon was passed through the vaporiser loaded with ZrCl4 powder so that it can carry the ZrCl4 vapour to the reaction chamber at the required flow and proportion. The stainless steel feed pipes carrying hydrogen, methane and ZrCl4 (under argon) were joined together and then fed in the reaction chamber. To avoid ZrCl4 agglomerating and clogging the inlet pipe, the feed pipes were heated to 300°C (ZrCl4 vaporiser temperature) by an electric heat tracing tape. Argon was also used as a purge gas and to stabilise the total reactor pressure.

! The control and transport of ZrCl4 vapour into the reaction chamber are crucial for the production of ZrC with the desired stoichiometry and morphology. ZrCl4 was first vaporised in the vaporisation system and the vapour was continuously swept to the reaction chamber using argon carrier gas. ZrCl4 was vaporised to the required vapour pressure in order to optimise the growth characteristics of the ZrC layers. The ZrCl4 vaporisation system consisted of an oven heated to 300°C (just below the phase transition temperature of 331°C)

Table II

:629854;2,6+%:7;+69:7861;3.:28/82698543 .:28/8269854 Material Purity Thermal conductivity (Wm-1 °C-1) Coefficient of thermal expansion (°C-1) Inner diameter (cm) Outer diameter (cm) Length (cm) Density (gcm-3) Grain size (Οm) Porosity

:38-4; 610: Graphite 100.00% 8.5 4.3Ă—10-6 2.5 3.5 30 1.77 15 9%

! !

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The gas delivery system comprised three separate gas cylinders (methane, argon and hydrogen cylinders) fitted with pressure gauges and flow meters. The precursors used in deposition of ZrC were anhydrous zirconium tetrachloride powder (purity > 99.5%; Hf < 50 mg kg-1 manufactured by Sigma-Aldrich (Pty) Ltd) and methane (CH4, 99.99% pure). ZrCl4 was carried from the vaporisation chamber to the reaction chamber by argon (Ar, 99.999% pure). Hydrogen gas (H2, 99.999% pure) was used to provide a reducing and diluting environment for the ZrCl4 vapour and the HCl effluent. Argon was also used to flush the reactor by continuously (four times before each deposition) pressurising and

Design and fabrication of a chemical vapour deposition system and a cylindrical steel vessel 24.6 cm long with inner diameter of 4.3 cm. The loaded cylindrical steel pot was connected to the inlet pipes and argon was allowed to sweep through freely. Several studies have been conducted on the CVD growth of ZrC from zirconium halides. It has been noted that ZrCl4, as opposed to other halides such as ZrBr4 and ZrI4, produces ZrC with relative good stoichiometry and is less toxic and relatively more volatile (Park et al., 2008). That is the reason it was selected as precursor. The physical properties of ZrCl4 are listed in Table III. Since ZrCl4 powder is very hygroscopic, precautions were taken to avoid its exposure to atmospheric moisture as far as possible. To determine and optimise the vapour pressure of ZrCl4 for the deposition of ZrC layers, ZrCl4 was heated under vacuum and under argon for 3 minutes for each temperature. Figure 3 shows a plot of the gauge pressure as a function of temperature, measured under vacuum and at argon gauge pressure of 11 kPa. For ZrCl4 under argon, the total pressure measured was the sum of the ZrCl4 (vapour) and argon partial pressures. It was observed that the presence of argon in the ZrCl4 does not change the pressure-temperature trend of ZrCl4 as shown in Figure 3. From room temperature up to 150°C the increase in pressure of ZrCl4 is small and then there is a rapid increase from around 190°C to about 330°C. A quick rise in vapour pressure was due to an increase in ZrCl4 vapour molecules as the temperature is increased (Liu et al., 2008). After 330°C the trend of the curve changed, indicating full sublimation of the ZrCl4. To avoid full sublimation and keep the vapour pressure at a manageable level, a temperature of 300°C was chosen throughout this study. We also investigated, by thermodynamic analysis, the ZrCl4 vaporisation process. It was found that the vaporisation of ZrCl4 does not generate other halide phases or form other lower halides such as ZrCl3, ZrCl2 or ZrCl as illustrated in Figure 4. This is because the Gibbs free energies of formation of ZrCl3, ZrCl2 and ZrCl are all greater than zero.

The ZrCl4 mass transfer rate as a function of argon flow rate is shown in Figure 5. The calibration curve was obtained by flowing argon through the ZrCl4 powder at different argon flow rates each for 20 minutes and determining the mass difference every after each run. The mass transfer rate ZrCl4 for each run was determined from Equation [3], derived experimentally: [3] where Mb is the mass of the loaded ZrCl4 vaporiser before vaporisation and argon flow and Ma is the mass of the loaded ZrCl4 vaporiser after vaporisation and argon flow for time t.

! The exhaust system was another essential part of the system. It consisted of a vacuum pump, cooling water jacket, scrubber containing calcium carbonate chips (CaCO3) and extraction ventilation. The hot exit gases from the reaction chamber were first passed through a coiled steel pipe equipped with a cooling water jacket before entering the scrubber containing calcium carbonate. The reason for passing the exit gases through the calcium carbonate scrubber was to remove hydrogen chloride gas and any unreacted ZrCl4 before it was released through the extraction system. Within the scrubber, a dust trap was installed to

! ! ! As mentioned, the ZrCl4 vapour was delivered into the reaction chamber by an argon flow, subsequently mixed with the methane and the hydrogen feed. The vapour pressure and the mass transfer rate of ZrCl4 depend mainly on the temperature at which ZrCl4 is vaporised and the argon flow rate. The mass flow rate of ZrCl4 was controlled by varying the argon flow rate and keeping the temperature at 300°C.

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Table III

,&38261; 75.:798:3;5/; 7(1 69;69+53.,:782 .7:3307:;" 69468 ; )) ; 04- ! ; )' # .:28/8269854

75.:79& Molecular weight Appearance State at room temperature Boiling temperature Melting temperature Latent heat of sublimation Specific gravity

233.03 g/mol at 15°C White Powder 331°C 437°C 4.54Ă—105 J/kg at 331°C 2.805 g/cm3 at 15°C

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Design and fabrication of a chemical vapour deposition system filter the exhaust gases before release into the environment. To regulate the reactor total pressure and to avoid accumulation of the exit gases in the reactor, which may otherwise cause contamination of the deposit, the reactor system was connected to a rotary vane vacuum pump. Hydrogen and methane are explosive and flammable gases, whose flammabilities are enhanced by high temperature, pressure and the presence of oxygen. Since the experimental environment was characterised by these conditions, caution was employed to reduce the associated dangers. The maximum flow rates of hydrogen and methane commensurate with safety and environmental considerations were determined. This was based on the minimum concentrations of hydrogen and methane necessary to support combustion, referred to as lower flammable or explosive limit (LFL or LEL). The standard LFLs of hydrogen and methane at room temperatures are 4% and 5% by volume respectively. The maximum hydrogen or methane flow rates that were used for these experiments were 42 cm3/min (approx. 0.7 cm3/s) and 870 cm3/min. These values were assumed to be the maximum leakage rates. Therefore, the percentage hydrogen or methane concentrations in the extraction were found to be 0.070% and 0.0034%, which were below the LFLs of 4% and 5% for hydrogen and methane respectively. Therefore the Necsa restriction of the hydrogen and methane concentrations being below 20% of its LFL was also satisfied. Before and during each experiment the gas leakages along the gas supply lines, the reactor system and the exhaust lines were monitored. This was done by (1) pressurising the reactor with argon and leaving it for over two hours and monitoring the pressure drop, (2) checking the joints with soap bubbles and (3) using a hydrogen leak detector. When the required minimum leakage rate was achieved, the CVD system was considered fit for deposition. In fact whenever there was a leak, the ZrC layer hardly grew on the substrate and when any deposition did occur the layer uniformity was compromised.

acetone, ethanol and then demineralised water for 20 minutes each. The cleaning and polishing was done in order to remove any oils, oxides and other contaminants so as to enhance adhesion and reduce the number of impurities in the layer being deposited. The clean substrates were dried in an oven at 200°C for 2 hours and then weighed to determine the mass of each substrate before deposition. The substrates were then kept in airtight desiccators ready for deposition. Prior to deposition, the substrates were mounted on the clean graphite substrate stage inside the graphite reaction chamber. The substrate stage was cleaned in a similar way as the substrates themselves to avoid contamination. The substrate stage was positioned such that its top, where the substrates sit, was within the area circumscribed by the induction coil. This was done to maximise and control the temperature.

//:29;5/;30%39769: 841:9;-6.;64*;-63; :152898:3;54;9,: -63;/15 ;*&46+823;84;9,:;7:62957 In an effort to understand and explain the complex gas flow mechanisms and layer growth characteristics, the finite element analysis software, COMSOL MultiphysicsÂŽ 5.1 was used to study the temperature and velocity profiles in the reactor. The effects of substrate-in let gap and gas flow velocities on the flow dynamics were simulated (Figure 6 and Figure 7). This numerical model analysis incorporates the conservation equations of momentum, energy and species. The model simulation results revealed that approximately uniform temperature and velocity profiles existed in the reactor, the magnitudes of which depend on the substrateinlet gap and gas velocity. The results of this simulation are important for the optimisation of the CVD operating conditions to achieve a high and uniform ZrC growth in the vertical reactor. This was verified by calculating the boundary layer thickness. The boundary layer thickness was calculated from Equation [4] (Pierson, 1999):

! The substrates of average diameter 10 mm and thickness 2.5 mm were cut from the bulk high-density (1.71g cm-3) graphite discs. These substrates had the same physical properties as those of the graphite reaction chamber listed in Table II. These graphite substrates were hand polished on a polishing wheel using 1000 grit silicon carbide paper. They were then sequentially cleaned by ultrasonic agitation with

[4] Here Ρ is the viscosity of the reactants, Ď is the density, V is the velocity of reactants, Re is the Reynolds number and d is the diameter of the reaction chamber. The detailed calculations of these variables plotted in Figure 8a and 8b are given in Biira et al. (2017a).

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Design and fabrication of a chemical vapour deposition system

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Keeping other factors constant it was found that increasing the substrate-gas inlet gap increased the boundary layer thickness and increasing the gas flow velocity reduces the boundary layer thickness. The higher the boundary layer thickness the thinner the layer thickness because fewer reactants reach the substrate. This in turn also affects the quality and uniformity of the growing surfaces.

//:29;5/;*:.5389854;9:+.:76907:3;54;9,:;-75 9, 2,67629:7839823;5/; (;16&:73 ! ! ! ! The average growth rate of the ZrC layers was calculated from the mass gain, the surface area of the substrate and theoretical density of ZrC (6.59 g/cm3) (Biira et al., 2017b). The deposition rate of the ZrC layers increased with increase in substrate temperature as shown in Figure 9a. There was a sharp increase in growth rate between 1200°C and 1300°C and a more gradual increase thereafter. This is because increasing temperature increases both the rate of decomposition of the reacting species and their reaction kinetics at the substrate surface. The linear fit of the Arrhenius plot in Figure 9b gives an apparent activation energy of 38.66 kJ/mol. However, it looks as if there might be a two-regime growth mechanism; one between 1200°C and 1300°C with an apparent activation energy value of 97.87 kJ/mol and the other between 1300°C and 1600°C with an apparent activation energy value of 19.28 kJ/mol. The apparent activation energy of the ZrC CVD reaction was determined from the Arrhenius Law, as given in Equation [5] (Park and Sudarshan, 2001):

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[5] where Ea is the apparent activation energy, T is the deposition temperature; R is the gas constant and A is the constant. The growth rate, k, was calculated as described in Biira et al., 2017b). Apparent activation energy, which is a function of reactant concentration and temperature, can be used to determine the rate limiting growth mechanism (Kashani, Sohi and Kaypour, 2000; Wagner, Mitterer and Penoy, 2008). Therefore the two regimes on the Arrhenius plot might be a representation of surface reactions as the limiting mechanism between 1200°C and 1300°C, since apparent activation energy is high, nd mass transport mechanism for the substrate temperatures between 1300°C and 1600°C for the low apparent activation energy of 19.28 kJ/mol.

! ! ! ! The structure of the as-deposited ZrC layers was analysed using a Bruker XRD D8 Advance with a Cu KÎą radiation source ( Îť=1.5406 Ă…). The working potential and current were set at 40 kV and 40 mA respectively. Figure 10 shows the XRD patterns of ZrC layers for substrate temperature ranging from 1200°C to 1600°C. At 1200°C only ten reflections of the ZrC phase were observed, viz. (111), (200), (220), (311), (222), (400), (331), (420), (422) and (511). These plane reflections indicated that the polycrystalline facecentred cubic structure of the ZrC layer was deposited, when matched with International Centre for Diffraction Data (ICDD) file number 03-065-8833. As the temperature was increased

Design and fabrication of a chemical vapour deposition system

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ZrC layers may partly originate from the presence carbon impurities in the lattice, the amount of which increased as the substrate temperature was increased, as shown in Figure 10. To confirm this, we added excess carbon onto the stoichiometric ZrC structure to evaluate some possible structures. This was done by the help of an electron structure software program for solids, surfaces and interfaces; the Vienna Ab initio Simulation Package (VASP) V4.6. It was observed that the average lattice parameter value increased as excess carbon is added, as illustrated in Figure 11. The presence of carbon impurities decreases the number of neighbour atoms of the Zr-C structure. This gives rise to relaxation of the lattice, causing the lattice parameter to increase. $8-07:;')! ;.699:743;/57; 7(;*:.5389:*;69;*8//:7:49;9:+.:76907:3

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from 1300°C to 1600°C some carbon peaks started to emerge, the intensities of which became increasingly prominent as the substrate temperature increased. The presence of free carbon in ZrC with increase deposition temperatures was also reported by Wang et al. ( 2008). The increase in carbon content may be attributed to the differences in the decomposition rates of methane and ZrCl4 as the temperature in the reaction zone is increased. This means that there is more atomic carbon than Zr and the excess carbon is deposited as free carbon with increasing temperature. The lattice parameters of the ZrC layers at different temperatures were determined from the Miller indices (hkl) and interplanar spacing of the various planes (Cullity and Stock, 1956; Nelson and Riley, 1945). The lattice parameter was found to increase from 4.6763 Ă… to 4.7011 Ă…. The change in the lattice constant values of the as-prepared ZrC layers may be due to strain in the layers. The strain in the

The surface morphology of the as-deposited ZrC layers was characterised by field emission scanning electron microscopy (FE-SEM) using a Zeiss Ultra Plus instrument at an acceleration voltage of 1 kV. Figure 12 shows the surface morphology of ZrC coatings obtained at 1200°C to 1600°C. At 1200°C the coating ball-shaped grains cling together forming cauliflower-like grains surrounded by abundant pores. This indicates that small grains agglomerate to form much larger grains. The ball-shaped grains become much dense and broader as the temperature is increased from 1300°C to 1600°C. No obvious cracks were observed at all substrate temperatures. The reduction of the numerous pores as the temperature is increased may be due to the fact that increasing temperature increases the energy of the reacting species, which may enhance surface diffusion at the substrate surface. As the temperature is increased, more of the reacting species are decomposed and therefore readily combine to form the layer. This increases the size of the growing crystals. It be can be concluded that the substrate temperature controls the morphology of ZrC coatings and the uniformity of the coating growth (Li et al., 2009).

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Design and fabrication of a chemical vapour deposition system

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tetrachloride purification process. International Journal of Chemical Engineering and Applications, vol. 3, no. 6. pp. 427–429.

(542103854 An inexpensive thermal CVD system was successfully designed and fabricated in-house for depositing a variety of films and coatings for various research and industrial applications. System parameters were optimised so as to achieve the required coating. These parameters included temperature, substrate-inlet gap and gas flow velocity. To confirm the capability of the CVD system, ZrC layers were successfully deposited on graphite substrates at atmospheric pressure for 2 hours at different temperatures. The results indicate that the substrate temperature has a determining influence on growth kinetics, surface morphology and microstructure of ZrC layers. Generally, increasing the substrate temperature increases the deposition rate of ZrC layers. Surface reaction is the controlling mechanism between 1200°C and 1300°C, whereas mass transport controls the deposition process between 1300°C and 1600°C. The results also show that the lattice parameter and the carbon content in the ZrC layers increase with substrate temperature. SEM micrographs show that increasing temperature increases the size of the particles, with the layer surface becoming more uniform with no noticeable cavities.

Necsa and the Department of Science and Technology of South Africa through the Nuclear Materials Development Network of the Advanced Metals Initiative are gratefully thanked for the provision of experimental materials and laboratory space. S. Biira acknowledges the financial support from the University of Pretoria and Busitema University.

:/:7:42:3 BIIRA, S., ALAWAD, B.A.B., HLATSHWAYO, T. T., CROUSE, P.L., MALHERBE, J.B., BISSETT, H., NTSOANE, T.P. and NEL, J.T. 2017a. Influence of the substrate gas-inlet gap on the growth rate, morphology and microstructure of zirconium carbide films grown by chemical vapour deposition. Ceramics International, vol. 43, no. 1. pp. 1354–1361. BIIRA, S., CROUSE, P.L., BISSET, H., ALAWAD, B.A.B., HLATSHWAYO, T.T., NEL, J.T. and MALHERBE, J.B. 2017b. Optimisation of the synthesis of ZrC coatings in an RF induction-heating CVD system using surface response methodology. Thin Solid Films, vol. 624. pp. 61–69. CULLITY, B.D. and STOCK, S.R. 1956. Elements of X-ray Diffraction. 1st edn. Cohen, M. (ed.). Addison-Wesley, Massachusetts. JUNG, Y.M., GYEONG, M.K. and MOONYONG, L. 2012. Measurement of bubble point pressures of zirconium and hafnium tetrachloride mixture for zirconium

KATOH, Y., VASUDEVAMURTHY, G., TAKASHI, N. and SNEAD, L.L. 2013. Properties of zirconium carbide for nuclear fuel applications. Journal of Nuclear Materials, vol. 441, nos.1–3. pp. 718–742. LI, H., ZHANG, L., CHENG, L. and WANG, Y. 2009. Oxidation analysis of 2D C/ZrC – SiC composites with different coating structures in CH4 combustion gas environment. Ceramics International, vol. 35. pp. 2277–2282. LIU, C., LIU, B., SHAO, Y.L., LI, Z.Q. and TANG, C.H. 2008. Vapor pressure and thermochemical properties of ZrCl4 for ZrC coating of coated fuel particles. Transactions of Nonferrous Metals Society of China (English Edition), vol. 18 (863614202), pp. 728–732. MEYER, M.K., FIELDING, R. and GAN, J. 2007. Fuel development for gas-cooled fast reactors. Journal of Nuclear Materials, vol. 371, no. 1. pp. 281–287. NELSON, J.B. and RILEY, D.P. 1945. An experimental investigation of extrapolation methods in the derivation of accurate unit-cell dimensions of crystals. Proceedings of the Physical Society, vol. 57, no. 3. pp. 160–177. PARK, J. and Sudarshan, T.S. 2001. Chemical Vapor Deposition. ASM International, Chicago. PARK, J.H., JUNG, C.H., KIM, D.J. and PARK, J.Y. 2008. Temperature dependency of the LPCVD growth of ZrC with the ZrCl4-CH4-H2 system. Surface and Coatings Technology, vol. 203, nos. 3–4. pp. 324–328. PATNAIK, P. 2003. Handbook of Inorganic Chemicals. McGraw-Hill, New York.

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KASHANI, H., SOHI, M.H. and KAYPOUR, H., 2000. Microstructural and physical properties of titanium nitride coatings produced by CVD process. Materials Science and Engineering A, vol. 286, no. 2. pp. 324–330.

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PIERSON, H.O. 1999. Handbook of Chemical Vapor Deposition: Principles, Technology and Applications. 2nd edn. William Andrew, New York. PIERSON, H.O. 1996) Handbook of Refractory Carbides and Nitrides: Properties, Characteristics, Processing and Applications. Noyes Publications, New Jersey. SIMPSON, P.G. 1960. Induction heating: coil and system design. McGraw-Hill, New York. WAGNER, J., MITTERER, C. and PENOY, M. 2008. The effect of deposition temperature on microstructure and properties of thermal CVD TiN coatings. International Journal of Refractory Metals and Hard Materials, vol. 26, no. 2. pp. 120–126. Wai, H.P., Aung Jr, S.S. and Win, T. 2008. Work coil design used in induction hardening machine. Work Academy of Science, Engineering and Technology, vol. 44. pp. 352–356. WANG, Y., LIU, Q., LIU, J., ZHANG, L. and CHENG, L. 2008. Deposition mechanism for chemical vapor deposition of zirconium carbide coatings. Journal of the American Ceramic Society, vol. 91 (23650). pp. 1249–1252. WON, Y.S., VARANASI, V.G., KRYLIOUK, O. ANDERSON, T.J., MCELWEE-WHITE, L. and PEREZ, R.J. 2007. Equilibrium analysis of zirconium carbide CVD growth. Journal of Crystal Growth, vol. 307, no. 2. pp. 302–308. YAN, X.T. and XU, Y. 2010. Chemical Vapour Deposition: an Integrated Engineering Design for Advanced Materials. 1st edn. Springer-Verlag, London.

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a3

Selective sublimation/desublimation separation of ZrF4 and HfF4 by C.J. Postma*, S.J. Lubbe†and P.L. Crouseâ€

The separation of zirconium and hafnium, which is essential in the nuclear industry, is difficult due to the great similarities in their chemical and physical properties. In contrast to the traditional aqueous chloride separation systems, the current process focuses on dry fluoride-based technologies, which produce much lower volumes of chemical waste. In the present work, separation is achieved in both a sublimation and a desublimation step, where the Zr/Hf mole ratio varies between 160 and 245 across the length of desublimer and 86 to 40 within the sublimer. Model predictions for the sublimation/desublimation rates fit the experimental results well, with deviations becoming more apparent as sublimation proceeds. This may be attributed to crust formation preventing the system from reaching thermodynamic equilibrium. The model adequately predicts time- and temperature-dependent mole ratios of both the sublimer residue and of the desublimed mass. 2,4 &* % zirconium, hafnium, fluoride, sublimation, desublimation.

1$)*& ( )+&$ Zr ores typically contain between 1 and 3 wt.% Hf (Xu et al., 2012). Zr metal for use in the nuclear industry is required to have a Hf content <100 ppm, owing to its high neutron cross-section (Brown and Healy, 1978). Therefore, the separation step is crucial in the preparation of nuclear-grade Zr metal, but this is considered to be very difficult due to the close similarities in the chemical properties of Zr and Hf (Smolik, Jakóbik-Kolon and Porański, 2009). Generally, the preparation of Hf-free Zr relies on the traditional wet routes, for example solvent extraction systems (Banda, Lee and Lee., 2012; Brown and Healy, 1978; Deorkar and Khopkar, 1991; Taghizadeh, Ghanadi and Zolfonoun, 2011, 2008; Xu et al., 2012; Yang, Fane and Pin, 2002). In contrast to the traditional aqueous chloride systems, the dry processes have the advantage of producing much less hazardous chemical waste. The New Metals Development Network (NMDN) under the Advanced Metals Initiative (AMI) was created by the Department of Science and Technology (DST) of South Africa in 2006, with the mandate to develop new processes for beneficiation relating to the

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* The South African Nuclear Energy Corporation SOC Ltd. (Necsa), Pretoria, South Africa. †Department of Chemical Engineering, University of Pretoria, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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4$& %+%

elements Zr, Hf, Ta and Nb. The current work demonstrates an attempt to develop a process involving the beneficiation of the mineral zircon (ZrSiO4) to produce nuclear-grade Zr metal. The zircon contains Hf as an impurity. To indicate this, the formula Zr(Hf)SiO4 is used throughout this text. In the zircon crystal structure, Zr and Hf occur in identical crystallographic positions. The proposed Zr metal recovery process involves plasma dissociation of the Zr(Hf)SiO4, which results in a much higher chemical reactivity. Here the zircon particles condense upon cooling as microcrystals of Zr(Hf)O2 embedded in an amorphous silica matrix, which is referred to as plasmadissociated zircon (PDZ, i.e. Zr(Hf)O2 SiO2). The PDZ is then desilicated (DPDZ or Zr(Hf)O2) and fluorinated with ammonium bifluoride (ABF) to produce Zr(Hf)F4. Separation of Hf from the Zr can be achieved using either the tetrafluoride or tetrachloride form, the latter of which will require an additional chlorination step. The next step is the plasma reduction of the tetrahalide to the metal and finally, purification of the metal powder by means of a high-temperature vacuum furnace (Nel et al., 2013; Retief et al., 2011). The nomenclature Zr(Hf)F4 defines a ZrF4 crystal structure in which 1 to 3% of the Hf is substituted for Zr. ZrF4 and HfF4 behave differently in the vapour phase due to the differences in vapour pressure. This implies that HfF4 is thermodynamically more stable than ZrF4 at a specific temperature. It is therefore assumed that the ZrF4 will sublime separately from the HfF4 due to the differences in thermodynamic stability of the two compounds.

Selective sublimation/desublimation separation of ZrF4 and HfF4 In this process, separation of ZrF4 and HfF4 is achieved using sublimation followed by desublimation. The separation involves the sublimation of the tetrafluorides in an inert atmosphere under controlled conditions. The sublimed mass (at approx. 800°C) diffuses into nitrogen, which is then passed across a water-cooled desublimer (annulus) with the aim of desubliming the one metal fluoride in preference to the other. This implies that separation is achieved in both the sublimer and desublimer, due to differences in both the sublimation and desublimation rates. These rates depend on the vapour pressures at the respective temperatures as well as the concentration and/or partial pressure differences. The aim is for the sublimer residue to be Hf-rich and the desublimer content Zr-rich. Zr/Hf content is determined by means of ICPOES analysis. The aim is to establish experimental conditions, i.e. sublimation time, temperature and position on the desublimer that provide optimal separation conditions. These conditions must, however, be compatible with economic considerations, as higher temperatures and longer sublimation runs will entail higher operating costs. This work is a continuation of previous work (Postma, Niemand and Crouse, 2015) with the same sublimation model but based on a different sublimation and desublimation (water-cooled) system. A desublimation model as well as experimental work is also included and compared to modelling results.

+),*')(*,-*, +, Sublimation is a general method used for the purification of ZrF4 by removing most trace elements, e.g. Fe, Co, Ni and Cu (Abate and Wilhelm, 1951; Dai et al., 1992; Kotsar’ et al., 2001; MacFarlane, Newman and Voelkel, 2002; Pastor and Robinson, 1986; Solov’ev and Malyutina, 2002a; Yeatts and Rainey, 1965). Sublimation methods for the separation of Zr and Hf have been reported in the literature, but these methods are all carried out under vacuum conditions (Monnahela et al., 2013; Solov’ev and Malyutina, 2002b). In the work done by MacFarlane, Newman and Voelkel (2002) the area-dependant rate of sublimation of ZrF4 was calculated and a value of approximately 1.87 g.m-2.s-1 was obtained at 850 to 875°C. Ti, Esyutin and Scherbinin (1990a, 1990b) found that pure ZrF4 has a higher sublimation rate than industrial-grade ZrF4, which contains a degree of impurities. They concluded that this phenomenon might be due to the accumulation of low-volatile components in the near-surface layer of the sample, making diffusion and

+ (*,- +),*')(*,- ' &(*- *,%%(*,%-!&*- * '$ - !

evaporation increasingly difficult, resulting in decreased sublimation flux. This led to the assumption that there might be a possibility of ‘crust-formation’ which could limit the sublimation rate to a certain extent. In a study on the influence of layer height on the vacuum sublimation rate of ZrF4, it was concluded that the sublimation rate does not necessarily depend on the height of the sample in the sublimator (Ti, Esyutin and Scherbinin, 1990c). This led to the assumption that the sublimation rates are area-dependent only and sublimation does not take place from the bulk of the material. Figure 1 gives a range of vapour pressures obtained from the literature for both ZrF4 and HfF4 at temperatures above 600°C (Benedict Pigford and Levi, 1981; Cantor et al., 1958; Koreneo et al., 1972; Sense et al., 1953, 1954).

3*& ,%%- ,% *+ )+&$ The block flow diagram of the process is shown in Figure 2. The sublimation rates of ZrF4 and HfF4 are strongly dependent on the vapour pressures and the mole fraction of the respective components in the bulk mixture and therefore the sublimation rate of the ZrF4 is much higher than that of the HfF4. The aim is to exploit the difference in vapour pressure between the two compounds for separation and to sublime a mass of Zr(Hf)F4 at a predetermined temperature for a predetermined time and to stop sublimation after the time has elapsed. The sublimed mass enters a desublimation zone in which the ZrF4 and HfF4 desublime at different rates. Separation is thus achieved in both the sublimer and the desublimer.

+ (*,- #& -!#& - +' *'"

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Selective sublimation/desublimation separation of ZrF4 and HfF4

+ (*,- ,"')+ -&!-%( #+",*-'$ - ,%( #+",*-+$%+ ,-) ,-)( ,-!(*$' ,

0 ,*+",$)'#- *& , (*, The Zr(Hf)F4 used in the experments was prepared by reacting Zr(Hf)O2 (originating from zircon) with ABF. The tube oven was heated to the set temperature (between 700 and 850°C) before sample loading. Nitrogen was used as the carrier gas throughout all experiments at a mass flow rate of 0.3 kg/h. Two types of experiment were undertaken. The first set of experiments was performed in order to determine the rate of sublimation, including the extent of separation of the Zr and Hf in the sublimer residue. To this end, several sublimers, each containing 15 g sample mass, were loaded into the tube furnace for a pre-determined time, each sublimer being individually loaded one after the other at

+ (*,- ,"')+ -&!-) ,-%( #+",*-(%,

different residence times.The weight losses (mass sublimed) as well as the Zr/Hf mole ratios were determined from the sublimer residue. No desublimer was used in these experiments for the determination of the rate of sublimation. In the second set of experiments, three sublimers each containing 15 g Zr(Hf)F4 were loaded in the tube furnace for 30 minutes. The aim was to desublime as much as possible on the length of the desublimer to determine the Zr/Hf mole ratio as a function of the desublimer length. It is assumed that the sublimation conditions are the same for the respective sublimer samples. The Zr/Hf mole ratios were determined by dissolving the Zr(Hf)F4 mixture in HF and determining the Zr and Hf contents by means of ICP-OES analysis.

& ,##+$ The diffusion coefficient can be estimated using the LennardJones potential to evaluate the influence of the molecular forces between the molecules. This correlation (Equation [1]), also known as the Chapman-Enskog equation, holds for binary gas mixtures of nonpolar, nonreacting species (Perry and Green, 1997; Welty, 2001), which is the case for ZrF4 and HfF4 in nitrogen. [1] where AB is the collision diameter, a Lennard-Jones parameter in Ă…, where A refers to nitrogen and B to either ZrF4 or HfF4. Since is denoted as the Lennard-Jones diameter of the respective spherical molecule (Welty, 2001), an estimation was made for the diameter of a ZrF4 and HfF4 molecule assuming sphericity. The sizes of the respective molecules were calculated at room temperature with the use of SpartanTM software (‘14 V1.1.4). The equilibrium geometry was calculated using the Hartree-Fock method with the 6-31* basis set. Estimated values for the collision diameters of ZrF4 and HfF4 with N2 were calculated as 4.205 and 4.185 Ă…, respectively. The collision integral ( D) is a dimensionless parameter and a function of the Boltzmann constant ( ), the temperature and the energy of molecular interaction AB, where AB= A B. The boiling points (Tb) for ZrF4 (912°C) and HfF4 (970°C) (Lide, 2007) were used to calculate the values for B (molecular interaction of ZrF4 or HfF4) with the use of an empirical correlation, given by Equation [2]: VOLUME 117

941

The sublimer residue as well as the desublimed mass was collected, weighed and analysed for Zr and Hf content to thereby establish the amount of sublimation/desublimation steps required to achieve nuclear grade purity. Figure 3 gives a schematic of the sublimer and desublimer inside the tube furnace. The section above the sublimation pan facilitates the flow of nitrogen gas, which reduces the partial pressure of the ZrF4 and HfF4 and carries the tetrafluorides to the desublimer space. The sublimer is a reactor boat (Figure 4) which is 100 mm long and 20 mm deep and can take a maximum of 80 g Zr(Hf)F4. The Zr(Hf)F4 to be sublimed is placed in the boat, which is positioned in the centre of the heating zone within the tube furnace. The desublimer is a long cylindrical pipe (cooled to approx. 30°C) inside another insulated pipe which facilitates the carrying of the gas mixture. A simple geometry is selected for the desublimer to facilitate easy removal of the components at the end of the experiment. The desublimer temperature is much lower than that of the outer pipe walls, in order to prevent desublimation on the outer pipe wall. The spacing between the two pipes is minimal to reduce the diffusion path of the condensing particles.

Selective sublimation/desublimation separation of ZrF4 and HfF4 [2] Estimated values for the energy of molecular interaction for ZrF4 and HfF4 in N2 were calculated as 4.305 Ă— 10-14 and 4.409 Ă— 10-14 erg respectively.

Kim (2013) demonstrated that the theoretical foundation of thermal diffusion relates to Einstein’s random walk theory and added the spatial heterogeneity of the ‘random walk’ to reflect the temperature gradient of thermal diffusion. The thermal diffusivity was then determined theoretically. The walk speed S corresponds to the speed of the Brownian particles, which is a function of temperature. The molecular description of the thermal diffusion coefficient is given by:

[8]

The net rate of mass efflux from the control volume therefore reduces to: [9] In order to calculate the change in total flux along the length of the sublimation pan, the pan is divided into segments of length Δz and the flux in each successive segment is calculated by adding the flux in the previous segment to the sublimed masses of ZrF4 and HfF4 in segment j of length Δz. The continuity equation therefore becomes: [10]

[3] which through integration yields (Nakashima and Takeyama, 1989; Kim, 2013):

where m ̇ i,j+1,tg is the total mass flow at time t of species i in the gas phase at the j th position.

[4]

The continuity equation in differential form is given by: [5] where is the moles per unit volume (or area), j=(Ď uĚ…) is the flux, where uĚ… is the flow velocity vector field. S is the ‘source’ or ‘sink’ term, which is the generation per unit volume (or area) per unit time. In this case the source term is the sublimation rate (ri). The system is operated in such a way that no accumulation of mass occurs within the control volume dĎ =0), which implies that the net rate of mass efflux from (— dt the control volume equals the rate of sublimation. The continuity equation reduces to: [6] where i denotes either of the two species ZrF4 and HfF4 and ri is the sublimation rate of ZrF4 (or HfF4) in moles per unit sublimation area per unit time. The rate model for the sublimation of ZrF4 and HfF4 is based on the work of Smith (2001), who predicted evaporation rates for liquid spills of chemical mixtures by employing vapour-liquid equilibrium. The sublimation rate is given by: [7] where ki is the mass transfer coefficient in m.s-1 at time t, Pi* is the vapour pressure in kPa, pi' is the partial pressure in the bulk gas, xi is the mole fraction of ZrF4 (or HfF4) in the unsublimed bulk mass, R is the ideal gas constant (8.314 Pa.m3. mol-1. -1) and T is the temperature in . Assuming that the flowing gas above the bed is homogenously mixed within each control volume, the net rate of mass efflux from the control volume can be written as follows:

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There are three ‘mechanical driving forces’ that tend to produce movement of a species with respect to the mean fluid motion. These are concentration gradient, pressure gradient and the external forces acting on the species. For purposes of this study, the assumption will be made that the effects of pressure gradient and external forces are negligible (Bird, Stewart and Lightfoot, 1960: 564-565). Therefore the diffusion currents consists of both density (Fick’s Law) and temperature gradients (Ludwig-Soret effect). This macroscopic description of thermodiffusion dates back to the 19th century and was developed in the context of standard kinetic theory and therefore applies to dilute gas mixtures only (Debbasch and Rivet, 2011). The mass flux can therefore be defined by two terms: (a) that of the concentration gradient, i.e. the concentration contribution to mass flux: [11] and (b) the temperature gradient, i.e. the thermal diffusion contribution to mass flux: [12] Di,N2 (previously defined as DAB) is the mass diffusion constant (m2.s-1) of the diffusing species i in nitrogen, Ď is the density of the bulk gas (kg.m-3), i is the mass fraction of the diffusion species, DiT is the thermal diffusion coefficient (a measure of the mass diffusion process due to temperature) and T the temperature (Bird et al., 1960: 502; Dominguez et al., 2011; Kim, 2013). The thermal diffusion term ( JiT) describes the tendency for species to diffuse under the influence of a temperature gradient. The species move toward colder regions and a concentration gradient is formed. The effect is small, but separations of mixtures can be effected with steep temperature gradients (Bird, Stewart and Lightfoot, 1960: 567; Kim, 2013). The total mass flux equation can now be written as:

Selective sublimation/desublimation separation of ZrF4 and HfF4 [13]

[14] The same applies for the temperature change, where it is assumed that The total mass flux equation now reduces to: [15] The concentration contribution to mass flux ( JA ) can therefore be written as:

+ (*,- ( #+"')+&$-*'),-&!- * ! '%-'-!($ )+&$-&!-) ,-%( #+"')+&$ )," ,*')(*,

[16] where ki is the mass transfer coefficient in the desublimer which is a function of the mass diffusivity and the film thickness, p'i,b is the partial pressure of species A in the bulk gas, p*i,w is the vapour pressure at the wall temperature and Tf is the film temperature, which is the average of the bulk and wall temperature. The temperature contribution to mass flux can therefore also be written as a function of the mass transfer coefficient:

+ (*,- ( #+",*-*,%+ (,-+$ + ')+$ - ' ,-!&*"')+&$

[17] In order to calculate the change in total flux along the length of the desublimer (annulus), the length is divided into segments of length Δz and the flux in each following segment is calculated by subtracting the desublimed masses of ZrF4 and HfF4 in the segment from the flux in the previous segment. The incremental mass balance in the desublimer along the length of the annulus can therefore be expressed by the following equation:

[18] + (*,- & ,##, -%( #+"')+&$-*'),%- &" '*, - +) -, ,*+",$)'# *,%(#)%

The weight fraction sublimed at different temperatures are reported in Figure 5. These rates are calculated based on the sublimer residue, with zircon content subtracted. One explanation for the sublimation rate forming a plateau before complete sublimation has been achieved (Figure 5) can be the formation of a crust-like surface or sintered cake (Figure 6) preventing further sublimation from occurring. This may be due to the presence of impurities originating from the zircon. The sublimation temperature stays the same, but sintering of the cake negatively influences the sublimation kinetics. The modelled rates are shown in Figure 7 and compared to the linear sections of the sublimation rates obtained

experimentally (Figure 5). The model accurately predicts the rates of sublimation at the lower temperatures, but seems to over-predict at the higher temperature of 790°C by a factor of 1.4. This over-prediction of the sublimation rate at 790°C may be even greater at higher temperatures. For instance, the model predicts an average total flux of 7.34 g.m-2.s-1 for ZrF4 and HfF4 subliming at 850°C. This predicted value is 3.9 times higher than the 1.87 g.m-2.s-1 value estimated from results by MacFarlane, Newman and Voelkel (2002). One reason for the area-dependent rates differing might be the presence of impurities in the sample, since this can have a direct influence on the rate of sublimation (Ti, Esyutin and Scherbinin, 1990b). With the model, the effect of impurities on the rate is not taken into account, which will result in a higher sublimation rate. VOLUME 117

943

5,%(#)%-'$ - +% (%%+&$

Selective sublimation/desublimation separation of ZrF4 and HfF4

Zr/Hf mole ratios for the sublimer residue ranged between 86:1 (starting material) and 30:1, depending on the sublimation temperature and duration. From Figure 8 it is evident that the sublimer residue becomes Hf-rich with time because HfF4 sublimes at a lower rate than ZrF4, which is mainly due to the lower sublimation temperature of the HfF4. Comparison of the linear section of the experimental Zr/Hf mole ratio data with that obtained from the model indicates that the model predicts too high a sublimation rate for HfF4, resulting in lower Zr/Hf mole ratios in the sublimer residue. One explanation may be that the vapour pressure data for HfF4 (only one source found in literature) is possibly wrong, which is likely due to the strong dependency of the vapour pressure on temperature, especially at high temperatures. The vapour pressure data for ZrF4 was taken from four sources with a standard error of 48.2 and 0.055 for the constants A and B in the vapour pressure correlation.

It is assumed that the sublimed ZrF4 and HfF4 desublimes into a crystal structure in which the Zr and Hf occur in identical crystallographic positions, as with the starting material. The mass desublimed on the desublimer is sampled at several intervals to determine whether the Zr/Hf ratio changes along the length of the desublimer. Figure 9 gives an illustration of the desublimer with desublimed Zr(Hf)F4 and the respective sampling points. Here X0 is the tip of the desublimer closest to the heating zone and along the length of the desublimer are X1 through to X5 which are 100 mm apart, with X1 being the first 10 mm on the desublimer. The sublimation temperatures investigated included 700, 740 and 790°C. Figure 10a gives an indication of separations achieved along the length of the desublimer. Figure 10b gives the Zr/Hf mole ratio of the collective mass obtained on the desublimer throughout. It is clear that the sublimation temperature does have an effect on the separation achieved in the desublimer. Model predictions are not shown here, but indicate better separation with mole ratios typically higher by a factor of 1.5 at 700 and 740°C and 1.2 at 790°C. There is currently no explanation for this. A lower sublimation temperature will result in better separation. Separation is further achieved along the length of the desublimer, which may be helpful in the design of a large-scale desublimer with different product collection points along the length. At 700 and 740°C approximately 50% of the sublimed mass desublimed and 70% of the sublimed mass desublimed at 790°C. The reason for this is still unknown. One explanation might be the concentration in the bulk gas stream, which is higher at the higher temperature.

+ (*,- * !-"&#,-*')+&-&!-) ,-%( #+",*-*,%+ (,-'%-'-!($ )+&$-&!-)+", '$ -)," ,*')(*,

+ (*,- '" #,- &%+)+&$%-&$-) ,- ,%( #+",*

(a)

Work reported in this paper is based on only a first sublimation step. This implies that the mass collected from the desublimer was not sublimed to account for a second or third step. The masses collected are simply too small to enable us to do second, third and so forth sublimation steps. The results collected from this first sublimation step were compared to the model and the comparisons used to estimate the extent of separation for second, third etc. sublimation steps. This is therefore a theoretical study to determine the

(b)

+ (*,- ' - , '*')+&$-'#&$ -) ,-#,$ ) -&!-) ,- ,%( #+",* - - &##, )+ ,-"'%%-'%-'-!($ )+&$-&!-)," ,*')(*,

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Selective sublimation/desublimation separation of ZrF4 and HfF4 5,!,*,$ ,% ABATE, L.J. and WILHELM, H.A. 1951. Sublimation of zirconium tetrafluoride. No. ISC-151. US Atomic Energy Commission, Ames Laboratory. BANDA, R., LEE, H.Y. and LEE, M.S. 2012. Separation of Zr from Hf in hydrochloric acid solution using amine-based extractants. Industrial and Engineering Chemistry Research, vol. 51, no. 28. pp. 9652–9660. BENEDICT, M., PIGFORD, T.H. and LEVI, H.W. 1981. Nuclear Chemical Engineering. McGraw-Hill. BIRD, R.B., STEWART, W.E. and LIGHTFOOT, E.N. 1960. Transport Phenomena. Wiley.

number of sublimation steps required to achieve < 100 ppm Hf in the product. The model was run at a sublimation temperature of 790°C for several repetitions, each time entering a new value for the Hf concentration (i.e. HfF4) in the bulk material with the assumption of removing the sublimer after 30 minutes. The Hf content calculated after each step was multiplied by 1.2 to account for inaccuracies in the model at this temperature. The assumption of 1.2 was incorporated since the model predictions deviated by a factor of 1.2 from the desublimation experimental results (i.e. Zr/Hf mole ratios). The results obtained (Figure 11) indicated that eight steps are required to reduce the Hf concentration to less than 100 ppm. With the assumption of a 70% collection efficiency as found in the desublimation experimental results at 790°C, it is calculated that 15.5 kg would be required to produce 1 kg of nuclear-grade ZrF4. There are many other factors that might influence the results, for example the particle surface area before and after sublimation and the amount of oxyfluorides present after collection from the desublimer. Also note that lower sublimation temperatures give better separation, but require longer sublimation times to achieve the same amount of sublimed material.

/&$ #(%+&$% The sublimation and desublimation model sufficiently predicts the sublimation rates as well as the separation of the ZrF4 and HfF4 in both the sublimer and the desublimer. Optimal temperature selection for sublimation is imperative, since lower temperatures result in a lower sublimation rate but give better separation, whereas higher temperatures require shorter residence times but give slightly less separation of the two respective components. Based on several assumptions, the model was used with a sublimation temperature of 790°C and a residence time of 30 minutes. The model results indicated that eight steps are required to reduce the Hf concentration to less than 100 ppm.

. $& #, ,",$)% The authors acknowledge the AMI of the DST and Necsa for providing the necessary funding and support to complete this study.

CANTOR, S., NEWTON, R., GRIMES, W. and BLANKENSHIP, F. 1958. Vapor pressures and derived thermodynamic information for the system Rbf-ZrF4. Journal of Physical Chemistry, vol. 62, no. 1. pp. 96–99. DAI, G., HUANG, J., CHENG, J., ZHANG, C., DONG, G. and WANG, K. 1992. A new preparation route for high-purity ZrF4. Journal of Non-Crystalline Solids, vol. 140, nos. 1-3. pp. 229–232. DEBBASCH, F. and RIVET, J.P. 2011 .The Ludwig–Soret effect and stochastic processes. Journal of Chemical Thermodynamics, vol. 43, no. 3. pp. 300–306. DEORKAR, N.V. and KHOPKAR, S.M. 1991. Liquid-liquid extraction of zirconium from hafnium and other elements with dicyclohexyl-18-crown-6. Analytica Chimica Acta, vol. 245. pp. 27–33. DOMINGUEZ, G., WILKINS, G. and THIEMENS, M.H. 2011. The Soret effect and isotopic fractionation in high-temperature silicate melts. Nature, vol. 473 (7345). pp. 70–73. KIM, Y.J. 2013. Einstein’s random walk and thermal diffusion. https://arxiv.org/abs/1307.4460 [accessed 12 May 2016]. KORENEO, Y., SOROKIN, I., CHIRINA, N. and NOVOSELO, A.V. 1972. Vapor-pressure of hafnium tetrafluoride. Zhurnal Neorganicheskoj Khimii, vol. 17, no. 5. p. 1195. KOTSAR’, M.L., BATEEV, V.B., BASKOV, P.B., SAKHAROV, V.V., FEDOROV, V.D. and SHATALOV, V.V. 2001. Preparation of high-purity ZrF4 and HfF4 for optical fibers and radiation-resistant glasses. Inorganic Materials, vol. 37, no. 10). pp. 1085–1091. LIDE, D.R. 2007. CRC Handbook of Chemistry and Physics, 88th edn. CRC Press, London. MACFARLANE, D.R., NEWMAN, P.J. and VOELKEL, A. 2002 Methods of purification of zirconium tetrafluoride for fluorozirconate glass. Journal of the American Ceramic Society, vol. 85, no. 6. pp. 1610–1612. MONNAHELA, O.S., AUGUSTYN, W.G., NEL, J.T., PRETORIUS, C.J. and WAGENER, J.B. 2013. The vacuum sublimation separation of zirconium and hafnium tetrafluoride. Proceedings of the AMI Precious Metals 2013 Conference, Protea Hotel, Cape Town, 14-16 October 2013. Southern African Institute of Mining and Metallurgy, Johannesburg. NAKASHIMA, K. and TAKEYAMA, N. 1989. Thermal diffusion in a Lorentz gas. Journal of the Physical Society of Japan, vol. 58, no. 12. pp. 4352-4357. NEL, J.T., HAVENGA, J.L., DU PLESSIS, W. and LE ROUX, J.P. 2013. Treatment of chemical feedstocks. US patent 9468975 B2. PASTOR, R.C. and ROBINSON, M. 1986. Method for preparing ultra-pure zirconium and hafnium tetrafluorides. US patent 4578252 A. PERRY, R.H. and GREEN, D.W. 1997. Perry's Chemical Engineers' Handbook. 7th Edition, McGraw-Hill, New York. POSTMA, J.J., NIEMAND, H.F. and CROUSE P.L. 2015. A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form. Journal of the Southern African Institute of Mining and Metallurgy, vol. 115, no. 10. pp. 961-965. RETIEF, W.L., NEL, J.T., DU PLESSIS, W. and CROUSE, P.L. 2011. Treatment of zirconia-based material with ammonium bi-fluoride. US patent 8778291 B2. VOLUME 117

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+ (*,- & ,#-*,%(#)- +) - &$%, ()+ ,-%( #+"')+&$-*($%

BROWN, A.E.P. and HEALY, T.V. 1978. Separation of zirconium from hafnium in nitric acid solutions by solvent extraction using dibutyl butylphosphonate: Part 1. Chemistry of the separation. Hydrometallurgy, vol. 3, no. 3. pp. 265–274.

Selective sublimation/desublimation separation of ZrF4 and HfF4 SENSE, K.A., SNYDER, M.J. and CLEGG, J.W. 1953. Vapor pressures of beryllium fluoride and zirconium fluoride. US Atomic Energy Commission Technical Information Services, Oak Ridge, Tennessee. SENSE, K.A., SNYDER, M.J. and FILBERT, R.B.J. 1954. The vapor pressure of zirconium fluoride. Journal of Physical Chemistry, vol. 58, no. 11. pp. 995-996. SMITH, R.L. 2001. Predicting evaporation rates and times for spills of chemical mixtures. Annals of Occupational Hygiene, vol. 45. pp. 437–445. SMOLIK, M., JAKÓBIK-KOLON, A. and PORA SKI, M. 2009. Separation of zirconium and hafnium using Diphonix® chelating ion-exchange resin. Hydrometallurgy, vol. 95, nos. 3–4. pp. 350–353. SOLOV’EV, A.I. and MALYUTINA, V.M. 2002a. Metallurgy of less-common and precious metals. Production of metallurgical semiproduct from zircon concentrate for use in production of plastic metallic zirconium. Russian Journal of Non-Ferrous Metals, vol. 43, no. 9. pp. 9–13. SOLOV’EV, A.I. and MALYUTINA, V.M. 2002b. Production of metallic zirconium tetrafluoride purified from hafnium to reactor purity. Russian Journal of Non-Ferrous Metals, 43 (9). pp. 14–18. TAGHIZADEH, M., GHANADI, M. and ZOLFONOUN, E. 2011. Separation of zirconium and hafnium by solvent extraction using mixture of TBP and Cyanex 923. Journal of Nuclear Materials, vol. 412, no. 3. pp. 334–337. TAGHIZADEH, M., GHASEMZADEH, R., ASHRAFIZADEH, S.N., SABERYAN, K. and MARAGHEH, M.G. 2008. Determination of optimum process conditions for

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the extraction and separation of zirconium and hafnium by solvent extraction. Hydrometallurgy, vol. 90, nos. 2–4. pp. 115–120. TI, V.A., ESYUTIN, V.S. and SCHERBININ, V.P. 1990a. The dependence of the zirconium tetrafluoride sublimation rate upon the process pressure. Kompleksnoe Ispol'zovanie Mineral'nogo Syr'ya, vol. 10. pp. 61–63. TI, V.A., ESYUTIN, V.S. and SCHERBININ, V.P. 1990b. The dependence of the zirconium tetrafluoride sublimation rate in vacuum upon the process temperature and product composition. Kompleksnoe Ispol'zovanie Mineral'nogo Syr'ya, vol. 9. pp. 63–64. TI, V.A., ESYUTIN, V.S. and SCHERBININ, V.P. 1990c. The influence of the sample height on the zirconium tetrafluoride sublimation process. Kompleksnoe Ispol'zovanie Mineral'nogo Syr'ya, vol. 8. pp. 60–61. WELTY, J.R., WICKS, C.E., WILSON, R.E. and RORRER, G. 2001. Fundamentals of Momentum, Heat and Mass Transfer. Wiley, New York. XU, Z., WANG, L., WU, Y., CHI, R., ZHANG, L. and WU, M. 2012. Solvent extraction of hafnium from thiocyanic acid medium in DIBK-TBP mixed system. Transactions of Nonferrous Metals Society of China, vol. 22, no. 7. pp. 1760–1765. YEATTS, L.B. and RAINEY, W.T. 1965. Purification of zirconium tetrafluoride. Technical report ORNL-TM-1292, US Atomic Energy Commission. YANG, X.J., FANE, A.G. and PIN, C. 2002. Separation of zirconium and hafnium using hollow fibers: Part I. Supported liquid membranes. Chemical Engineering Journal, vol. 8, nos. 1–3. pp. 37–44.

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946

VOLUME 117

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a4

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb by S.S. Mahlalela and P.G.H. Pistorius

The current study is intended to characterise the complex microstructure that results from autogenous welding of Zr-2.5Nb and its influence on mechanical properties. Laser beam welding (LBW) and gas tungsten arc welding (GTAW) were performed on a 1.5 mm thick samples by applying different combinations of speed and power to vary the heat input. The effect of cooling rate on β-ι phase transformation temperature and the resulting microstructure was investigated using a Bahr dilatometer. In addition, mechanical properties were determined by tensile and hardness testing. The base metal microstructure consisted of a fine structure of Zralpha with a small fraction of Nb-beta phase. Both LBW- and GTAWwelded joints displayed similar martensitic microstructures, with the main difference being the weld geometry and strength. CB)29>8= Zr-2.5Nb, autogenous welding, microstructure, β-ι transformation.

<A>9865A?9< Zirconium is used to manufacture nuclear plant components, more specifically the pressure tube containing the UO2 fuel rods. The use of zirconium and its alloys in the nuclear industry is credited to Admiral Rickover of the US Navy, who oversaw the development of the Shippingport Atomic Power Station in Pennsylvania (Krishnan and Asundi, 1980). This was the world's first commercial pressurized water reactor (PWR) used for generating electricity. Zirconium alloys are preferred because of their low neutron absorption cross-section, good resistance to corrosion, high strength and good creep resistance (Douglass, 1971). The desire by the nuclear industry to prolong the service life of cladding tubes while improving power output by increasing the operation temperatures led to the development of zirconium-niobium alloy. This alloy was first used in the USSR, due to its superior properties compared to other popular alloys like Zr-Sn (Cox, 2004). Currently these pressure tubes operate at about 300°C with a coolant pressure of about 10 MPa in pressurized heavy water reactors. The pressure tubes are extruded at around 820 to 850°C, cold worked 20 to 30% and stress relieved (autoclaved) at 400°C. The autoclaving

VOLUME 117

* SAIW Centre for Welding Engineering, Department of Material Science and Metallurgical Engineering, University of Pretoria, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

947

%)<93=?=

treatment at 400°C and the reactor operation temperatures of 300°C eventually leads to a microstructure of elongated grains with equilibrium phases (ÎąZr and βNb) Zirconium-2.5Nb exhibits two equilibrium phases, namely alpha (ÎąZr) phase with hexagonal close-packed (hcp) crystal structure at room temperature and beta (ÎąZr) phase with body-centred cubic (bcc) structure at temperatures above about 870°C (Banerjee and Mukhopadhyay, 2010). In the equilibrium phase diagram, the single-phase Îą and singlephase β regions are separated by a two-phase region of (ÎąZr+βZr) between 602–871°C. Alloying elements have an effect on the alpha to beta transformation temperature (Lustman and Kerze, 1955). Beta-stabilising elements (such as Sn and Nb) lower this temperature, thereby stabilising the ÎąZr phase. Alphastabilising elements (such as H, N and O) raise the alpha to beta transformation temperature, stabilising the βZr phase. Niobium is a βZr phase stabiliser, but conventional Zr-Nb alloys contain some oxygen, which is an ÎąZr-phase stabiliser. The phase transformation of beta to alpha during thermal treatment of Zr-2.5Nb alloy results in mainly two microstructural morphologies, namely Widmanstätten and martensitic structures. A Widmanstätten structure can further manifest itself in two morphologies, called parallel plates and basketweave (Holt, 1969). The parallel plate structure consists of a number of long ÎąZr plates that precipitate and grow on the same habit plane from the grain boundary in one parent β grain. Basketweave structure, in

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb Table I

4B;?5@: 59;39=?A?9< 9/ > +@=B ;BA@: 6=B8 ?< 56>>B<A =A68) 133; 6<:B== <9AB8 9A4B>2?=B0 !:B;B<A= Zr705 Limits (max) ASTM B352

+ 12A 0

*B

%?

$:

2.58

1100

1360

27

66

80

113

58

2.4–2.8

1500

900–1500

100

80

120

270

75

contrast, is composed of relatively short ÎąZr plates that randomly precipitate on a number of planes within the same parent β grain. The latter structure is promoted by an increase in cooling rate and the presence of randomly dispersed second-phase particles in the parent β phase. There are non-equilibrium phases such as the martensite (ιʹ) phase (hcp and orthorhombic), the ω- phase and a large number of metastable intermetallic phases. The retention of a metastable Zr-rich β-phase during fast cooling has also been found in the binary Zr–2.5Nb system (Banerjee and Mukhopadhyay, 2010). Zirconium-niobium alloys mainly show martensite with an acicular morphology (Srivastava et al., 2000). This is characterised by a structure consisting of primary martensite plates cutting across the parent β grain. Secondary plates subsequently transform from the primary plates in different variants of orientation relation (Banerjee and Mukhopadhyay, 2010). The deformation mechanism of zirconium alloys (hcp crystal structure) is through both the slip system and the deformation twin system. This complex behaviour is due to the less ideal axial ratio (c/a) of the alpha phase (Cochrane, 2013). A number of factors determine which system is activated first and these include microstructure, texture, deformation temperature, strain mode and strain rate. At intermediate temperatures (RT), transition in the ratecontrolling deformation mode from slip to twinning has been observed to occur during the course of plastic deformation (Song and Gray, 1995). Reducing grain size has been shown to inhibit twinning deformation mode, while increasing strain rate promotes it. Welding of the pressure tubes is one of the critical steps during the manufacturing process. Zr-2.5Nb has shown good weldability but its susceptibility to hydrogen, nitrogen, oxygen and carbon contamination requires serious consideration, especially at temperatures above 370°C (Krueger, 2015). This is the reason why argon and helium are the only shielding gases recommended and thorough cleaning of the workpiece before welding is essential. Contamination causes embrittlement, which results in increase in hardness but loss of ductility. Adequately shielded welds should show bright silver weld beads. When contaminated, zirconium welds show a range of discolouration from dark straw to blue for light contamination to dark blue or grey for severe contamination (Krueger, 2015). In this study, autogenous welding was done where no filler metal is added and the fusion zone is formed by the melting and resolidification of the base metal. Laser beam welding (LBW) and gas tungsten arc welding (GTAW) were used. Laser beam welding allows precision, high-quality weld 948

#/

VOLUME 117

joints to be produced with a low heat input (HI). This results in a small heat-affected zone (HAZ), fast cooling rate with very little distortion and a high depth-to-width ratio. GTAW produces acceptable quality with wider welds due to an inherently higher heat input. The microstructural change that occurs due to thermal treatment can have a negative bearing on the performance of zirconium alloys (Rudling, 2007). The aim of this study is to characterise the complex microstructure that results from autogenous welding and compare the results of the two welding processes. In addition, tensile and micro-Vickers hardness tests were done to quantify the mechanical behaviour of the various weld beads.

! 3B>?;B<A@: 3>95B86>B Zr705 alloy sheet of 1.5 mm thickness supplied by ATI Specialty Alloys & Components was used for all experiments. The chemical composition of the alloy as provided by the supplier is shown in Table I. The sheet was sectioned into 100Ă—50 mm rectangular samples for autogenous bead on plate welding. The samples were cleaned with ethanol before welding to limit contamination

An IPG YLR fibre laser was used with maximum power output of 3 kW, an emission wavelength of 1.07 Οm and a spot focus diameter of 250 Οm with a focal length of 200 mm. The welding head was mounted on an articulated robotic arm. The samples were assembled on a welding table with a root purge incorporated. Helium was used for root purging and shielding while a trailing weld shoe using argon was utilised for the top side of the weld bead. Visually, most welds were free of colouration, which suggests adequate shielding, except for two high heat input welds which were dark straw to light blue. These colours indicate surface oxidation, which can be removed using a wire brush. Table II shows the laser beam welding parameters that were chosen to ensure full penetration on both processes. The heat input, which is critical in determining the cooling rate, was calculated using Equation [1] (Weman, 2003). [1] where Q is heat input (kJ/mm), V is voltage (V), I is current (A), ν is welding speed (mm/min) and Ρ is arc efficiency or power transfer efficiency. For laser beam welding, the arc efficiency was taken to be Ρ = 0.8 (Fuerschbach, 1996). The reported weld widths in Table II were the average width of the top and the root of the fusion zone.

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb Table II

@=B> 2B:8?<7 3@>@;BAB>= 1 0 %@;3:B :@+B:

92B> 1.0

.B:8?<7 =3BB8 1;;&;?<0

#B@A ?<36A 1( &;;0

.B:8 2?8A4 1;;0

L1

1000

4000

0.012

1.1

L2

1500

0.018

1.3

L3

3000

0.036

1.4

L4

1000

0.024

1.4

L5

1500

0.036

1.5

L6

3000

0.072

1.8

L7

600

0.029

1.7

2000

1000

L8

800

0.038

1.8

L9

1000

0.048

2.2

Range

0.012 to 0.072

DCEN polarity was used with an EWTh 2.4 mm diameter electrode on the semi-automated GTAW machine. The electrode tip was ground to a 45° angle and arc length of 2 mm. The welding arc was initiated through the touch-start procedure. Argon was used for purging and shielding through the backing plate, torch nozzle and trailing weld shoe. No discolouration was observed on the welds visually, which suggested adequate shielding. Table III shows the gas tungsten arc welding parameters. The heat input was calculated as noted for laser beam welding, except that the arc efficiency was taken as Ρ=0.4 (Smartt, Stewart and Einerson, 1985).

Continuous cooling transformation behaviour was studied using varying thermal cycles in the dilatometer. A Bahr dilatometer was used with helium gas for better cooling control. An S-type thermocouple was spot-welded on the samples. The samples were heated at a constant rate of 10°C/s to a peak temperature of 1050°C and held for 10 minutes before cooling at different rates of 0.5, 1, 10, 25, 50, 150, 300 and 600°C/s. The data obtained was used to plot

the continuous cooling curves and temperature transformation and time diagram. Dilatometry samples were subjected to metallography and hardness testing.

Samples were prepared for microstructure examination by first grinding with SiC paper to 1200 grit. Chemical polishing, which simultaneously remove scratches and etches, was then performed using 10HF–45HNO3–45H2O2 (vol.%) solution. The microstructure was characterised using stereoscopic, optical and scanning electron microscopy (SEM).

Hardness testing was performed using a micro-Vickers machine with a load of 500 g force and dwell time of 10 seconds. The welded samples were tested on the transverse section (top welded side) across the whole profile from base metal, heat affected zone, to the weld with the indents 0.5 mm apart. Five measurements were done on the dilatometry and as-received samples. To evaluate and compare the mechanical properties of the welds, sub-size tensile samples with 25 mm gauge length and 6 mm width were prepared according to ASTM E8/E8M.

Table III

,$. 3@>@;BAB>= 1Ρ 0 .B:8?<7 56>>B<A 1$0

.B:8?<7 -9:A@7B 1 0

.B:8?<7 =3BB8 1;;&;?<0 205.8

T1

54.7

9.42

T2

69.7

10.43

T3

79.6

11

T4

70

10.5

#B@A ?<36A 1( &;;0

301.2

.B:8 2?8A4 1;;0

0.060

4.2

0.085

6.1

0.102

6.8

0.059

4.3

T5

90

11.1

0.080

5.2

T6

104.6

11.45

0.095

7.4

T7

99.7

11.1

0.059

4.2

T8

119.6

12.5

0.079

5.5

Range

452.4

0.059 to 0.102

VOLUME 117

949

%@;3:B

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb The samples were machined parallel to the weld line. Conventionally, the tensile properties of a welded joint characterised using a tensile test coupon that is machined transverse to the weld (AWS B4). It will be shown later that the hardness of the weld metal and of the heat affected zone was higher than the hardness of the base metal. A transverse tensile test coupon would therefore have failed in the unaffected base metal and the tensile behaviour of the weld metal would not have been characterised. Given the coupon width (6 mm) and the width of the weld bead (Tables II and III), the tensile test results therefore reflect the combined mechanical behaviour of the weld metal and the heat affected zone. The samples were deformed with a constant cross-head speed of 2 mm/min and all tests were carried out at room

temperature. Five samples were tested in total, including two samples of laser beam welds, two of GTAW welds and one of as-received sheet. The welding of all tensile samples was performed to ensure heat input between 0.065–0.070 kJ/mm.

DB=6:A= @<8 8?=56==?9<

Microstructural examination of the dilatometry samples cooled at a rate below 50°C/s showed mainly a basketweave structure in the interior of the parent beta grain with a parallel alpha plate structure on the grain boundary. These parallel plates nucleate from the grain boundary and grow into the prior beta grain. The beta (βNb) niobium phase is

*?76>B "' 3A?5@: ?;@7B= 9/ " &= 599:?<7 >@AB 8?:@A9;BA>) =@;3:B ;?5>9=A>65A6>B 1@0 ? A6>B 9/ +@=(BA2B@-B 3@>@::B: 3:@AB @<8 +BA@ 1 +0 <?9+?6; 34@=B 1+0 @=(BA2B@-B @A 4?74 ;@7<?/?5@A?9<

*?76>B ' ?:@A9;BA>) ;?5>9=A>65A6>B @A -@>)?<7 599:?<7 >@AB= %! ?;@7B= @A " ;@7<?/?5@A?9<

950

VOLUME 117

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb previously published continuous cooling transformation diagram (Saibaba, 2012). The current study showed that, at a cooling rate of below 50 C/s, the transformation product was a mixture of basketweave, parallel plate and beta niobium (Figure 1 and Figure 2). At a cooling rate higher than 50°C/s, martensitic phase was the dominant transformation product (see, for example, Figure 2). The change in transformation product was not associated with an abrupt change in transformation start or transformation end temperatures (Figure 3). The reason for the gradual and consistent change in transformation temperatures, while the transformation products changed markedly, is not clear.

Both laser and GTAW processes display similar weld microstructure constituents and characteristics, as shown in Figure 4 (the L9 microstructure shows an uneven weld

Temperature °C

observed mainly on the grain boundary. Figures 1a and 1b shows the microstructure at 1°C/s cooling rate. As the cooling rate increases, intragranular alpha plates nucleate from the interior of the beta grain and become progressively finer in size. At high cooling rates (≼ 50°C/s) a martensitic reaction occur which results in acicular (ιʹ) plates forming. The change in microstructure with increasing cooling rate is shown in Figure 2. The temperature-time transformation curves during cooling are shown in Figure 3. The highest transformation temperature of βι( ÎąZr+ βZr) is 856°C at 0.5°C/s cooling rate. The transformation start temperature was measured as 775°C where (βιιʹ) bcc beta transforms to hcp alpha martensite at 50°C/s. The (βιιʹ) transformation start temperature decreased to 650°C at 600°C/s cooling rate. At higher cooling rates, the transformation start and transformation end temperatures generally decreased, consistent with a

*?76>B ',,, 8?@7>@; 8?:@A9;BA>) =@;3:B= 599:?<7 56>-B= ,4B >B8 @<8 +:@5( ;@>(B>= ?<8?5@AB A4B AB;3B>@A6>B= @A A4B =A@>A @<8 /?<?=4 9/ A>@<=/9>;@A?9< >B=3B5A?-B:)

VOLUME 117

951

*?76>B '%AB>B9=593B ?;@7B= =492?<7 2B:8?<7 ;?5>9=A>65A6>B= @<8 @>B :@=B> 2B:8= 24?:B ," @<8 , @>B ,$. 2B:8= ,4B )B::92 @>>92 ?<8?5@AB= A4B 2B:8?<7 8?>B5A?9< 9< @:: A4B 2B:8=

Weld Width (mm)

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb

*?76>B '!//B5A 9/ 4B@A ?<36A 9< A4B 2B:8 2?8A4

Hardness values for dilatometry at slow cooling rate averaged 210 Hv500 (mixture of basketweave and parallel plate structure) and increase steadily to 242 Hv500 at 25°C/s (fine basketweave microstructure). The hardness of martensitic structure ranged between 270 and 280 Hv500 for cooling rates of 50°C/s and above. Hardness and cooling rate have a positive linear relationship until the martensitic transformation, as seen in Figure 6. Laser welds have superior strength to GTAW welds, as shown by the hardness results in Table IV, where the results for three welds of each welding process are presented. These were chosen from the low, medium and high heat input welds. The laser weld and HAZ hardness are 20–25 Hv500 higher than for GTAW. The tensile curves in Figure 7 are typical in zirconium deformation at RT, where initially yielding is smooth and the flow curve gradually plateaus, leading to a point in which the rate of dynamic recovery is equal to the hardening rate. As explained in the introduction, deformation in hcp zirconium alloys is accommodated by the activation of either the slip or

Hardness micro-Vickers Hv500

surface due to unintended over-etching). The main difference is the much larger width size of GTAW welds and HAZ as compared to those of laser welds (take note of the varying magnification used between the microstructures). The difference in weld bead size and morphology between L3 and L9 is mainly attributed to the varying heat input from 0.012 to 0.072 kJ/mm. Similarly, the width of the GTAW welds increased with heat input (Figure 5). The base metal microstructure was a fine cold-worked structure of alpha (ιZr) phase with very small volume fraction of βNb phase. The heat affected zone has recrystallised fine equiaxed grains of alpha with beta phases on the grain boundaries. The grain size increases toward the fusion line due to grain growth. The weld microstructure is dependent on the heat input and welding speed, but dominated by β grains that have transformed to martensite. At low heat input the microstructure is characterised by equiaxed grains. As the heat input increases, the structure initially consists of large equiaxed grains and the morphology changes to columnar grains towards the centre. In GTAW welds with high heat input, the weld grains are mainly large and columnar due to the slow welding speed.

*?76>B ' ?:@A9;BA>) =@;3:B= =492?<7 A4B B//B5A 9/ 599:?<7 >@AB 9< 4@>8<B== @<8 A4B 54@<7B= ?< ;?5>9=A>65A6>B

952

VOLUME 117

Microstructural characterization of laser beam and gas tungsten arc welded zirconium-2.5Nb Table IV

$-B>@7B 4@>8<B== AB=A >B=6:A= 1#- 0

Laser weld

GTAW

Base metal

.B:8?<7 4B@A ?<36A 1( &;;0

#$

.B:8

0.012

265

280

0.038

260

279

0.072

261

276

0.059

246

255

0.085

242

249

0.102

243

248

-

205–220

behaved in a similar way to the dilatometry microstructural response to change in cooling rate Laser beam welding results in stronger welds due to the fast welding speed and lower heat inputs. Finer martensite from laser welds imparts higher strengths but reduced ductility For nuclear application, laser beam welding would be more desirable because of the high welding speed, smaller weld bead size and higher strength.

$5(<92:B87B;B<A= The CSIR is acknowledged, especially Mr Herman Burger for his assistance and advice on laser beam welding. NECSA and AMI are thanked for their continued support.

DB/B>B<5B= BANERJEE, S. and MUKHOPADHYAY, P. 2010. Phase Transformations: Examples from Titanium and Zirconium, vol. 12. Chapter 1. Elsevier.

Stress (MPa)

COCHRANE, C.J. 2013. Effect of process parameters on deformation of Zr2.5wt%Nballoy. Masters thesis, Queen’s University, Canada COX, B. 2004. Development of Zr-Nb alloys. ZIRAT-9 Special Topic Report on Corrosion of Zr-Nb Alloys. Advanced Nuclear Technology International, Sweden. DOUGLASS, D.L. 1971. The metallurgy of zirconium. Atomic Energy Review Supplement. FUERSCHBACH, P.W. 1996. Measurement and prediction of energy transfer efficiency in laser beam welding. Welding Journal, vol. 75. pp. 24–34. HOLT, R.A. 1969. The beta to alpha to phase transformation in Zircaloy-4. Journal of Nuclear Materials, vol. 3. pp. 322–334.

*?76>B '%A>B== =A>@?< 56>-B= AB=AB8 @A >99; AB;3B>@A6>B 2?A4 5>9== 4B@8 =3BB8 9/ ;;&;?<

KRISHNAN, R. and ASUNDI, M.K. 1981. Zirconium alloy in nuclear industry. Proceedings of the Indian Academy of Sciences Section C: Engineering Sciences, vol. 4, no. 1. pp. 41–56. KRUEGER, B.R. 2015. Reactive refractory and precious metals and alloys. Welding Handbook. American Welding Society. Chapter 7, pp. 449–482.

9<5:6=?9<= The microstructure morphology responded to heat input. Equiaxed grains of martensite were observed at low heat input, but columnar grains that grew in size developed as heat input increased There was good agreement between dilatometry and the weld microstructure constituents and strength results. Increased heat input resulted in a slower cooling rate and consequently a decrease in tensile strength and hardness. The weld microstructures also

LANCASTER, J.F. 1993. Metallurgy of Welding. 5th edn. Chapman & Hall. LUSTMAN, B. and KERZE, F. 1955. The Metallurgy of Zirconium, McGraw-Hill, New York. OSLON, D.L. (ed.). 1993. ASM handbook: welding, brazing and soldering, vol. 6. ASM International, West Conshohocken, PA. RUDLING, P., STRASSER, A. and GARZAROLLI, F. 2007. Welding of zirconium alloys. Special report, Advanced Nuclear Technology International. SAIBABA, N., JHA, S., TONPE, S., VAIBHAW, K., DESHMUKH, V., RAO, S.R., KRISHNA, K.M., NEOGY, S., SRIVASTAVA, D., DEY, G. and KULKARNI, R. 2012. Microstructural studies of heat treated Zr-2.5 Nb alloy for pressure tube applications. Zirconium in the Nuclear Industry: Proceedings of the 16th International Symposium. ASTM International, West Conshohocken, PA. SMARTT, H.B., STEWART, J.A. and EINERSON, C.J. 1985. Heat transfer in gas tungsten arc welding. Proceedings of the ASM International Welding Congress, Canada. ASM International, West Conshohocken, PA. SONG, S.G. and GRAY, G.T. 1995. Influence of temperature and strain rate on slip and twinning behaviour of Zr. Metallurgical and Materials Transactions, 26A. pp. 2665–2675. SRIVASTAVA, D., MUKHOPADHYAY, P., BANERJEE, S. and RANGANATHAN, S. 2000. Morphology and substructure of lath martensite in dilute Zr-Nb alloys, Materials Science and Engineering, vol. A288. pp. 101–110. WEMAN, K. 2003. Welding Processes Handbook. Woodhead Publishing, UK. WILLIAMS, C.D. and GILBERT, R.W. 1966. Tempered structures of a Zr-2.5 wt %Nb alloy. Journal of Nuclear Materials, vol. 18. pp.161–166. VOLUME 117

953

twinning systems. At room temperature, zirconium initially deforms dominantly through slip until a transition (critical) point is reached, where the onset of uniform twinning deformation occurs. The transition point can be clearly seen in the ‘no weld’ curve. After the onset of plastic deformation, some strain softening occurred (at an engineering strain of about 3.5%), followed by a constant flow stress (up to a strain of about 15%). The laser beam welds had a higher strength, due to the faster cooling rate and a much finer martensitic structure. The weld metal tensile strength exceeded the strength of the base metal (Figure 7) and exceeded the minimum tensile strength for this alloy in the annealed condition (450 MPa, ASTM B352).

Primary Processes The CSIR Light Metals (LM) Competence Area hosts the Light Metals Development Network (LMDN) and Titanium Centre of Competence (TiCoC). Both TiCoC and LMDN are supported by DST through AMI. This is in response to the mineral beneficiation plan of SA as the country is the second largest producer of Titanium minerals with a share of 23 % of the world deposits. The CSIR, various universities, science councils and private companies are collaborating in developing the technology building blocks and R&D platforms towards adding value to the countries’ mineral resources. CSIR has specialised facilities and capabilities for materials (metals and composites) design and selection, materials testing and characterisation, structural design and analysis and advanced capabilities in parts/components manufacturing, which benefit key stakeholders, such as universities, new enterprises, SMEs and large enterprises. The CSIR LM Team of highly experienced and qualified researchers is equipped with advanced infrastructure for innovation and manufacturing Human capital capabilities. development is a major focus and partnerships with other key players in the innovation ecosystem are encouraged.

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http://dx.doi.org/10.17159/2411-9717/2017/v117n10a5

Pd0.02Ce0.98O2- : a copper- and ligandfree quasi-heterogeneous catalyst for aquacatalytic Sonogashira crosscoupling reaction by P.P. Mpungose, N.I. Sehloko, T. Cwele, G.E.M. Maguire and H.B. Friedrich

A Pd0.02Ce0.98O2- solid solution oxide was synthesised in one step using a urea-assisted solution combustion method. XRD, ICP-OES, BET, XPS, SEM, EDX, TEM, TGA and Raman spectroscopy were used to study the structural and electronic properties of the as-prepared Pd0.02Ce0.98O2- catalyst. The catalyst testing results revealed that the Pd0.02Ce0.98O2- system performs as an efficient precatalyst for the Sonogashira crosscoupling reactions under copper- and ligand-free conditions. A wide range of iodoarenes were efficiently coupled to phenylacetylene with good to excellent isolated yields. A thorough investigation through a series of suitable experiments explicitly showed that the Sonogashira crosscoupling reaction is accomplished via a quasi-heterogeneous mechanism by the leached Pd(0) species. As a result, the Pd0.02Ce0.98O2- lost some activity upon recycling. EB4 A<8; Pd2+ substituted ceria, solution combustion, nanoparticles and Sonogashira coupling reactions.

?><A85:>@A? The Sonogashira cross-coupling reaction is a great procedure for C-C bond formation between vinyl or aryl halides and terminal alkynes (Sonogashira, Tohda and Hagihara, 1975; Takahashi et al., 1980; Sonogashira, 2002). Sonogashira, Tohda and Hagihara (1975) recognised that the addition of copper(I) iodide significantly accelerates the alkenylation reaction that was conventionally catalysed by palladium−phosphane complexes (Sonogashira, Tohda and Hagihara, 1975; Takahashi et al., 1980; Sonogashira, 2002). Their discovery engendered great interest since the products of this reaction are important intermediates in the synthesis of organic materials, pharmaceuticals and natural products (Sonogashira, Tohda and Hagihara, 1975; Takahashi et al., 1980; Sonogashira, 2002; BĂśhm and Herrmann, 2000; Cwik, Hell and Figueras, 2006; Hosseini-Sarvari, Razmi and Doroodmand, 2014; Mujahidin and Doye, 2005; King and Yasuda, 2004). For example, the Sonogashira coupling procedure is used in the synthesis of the drug Terbinafine, which is used as an antifungal agent (Torborg and Beller, 2009; Beutler et al., 1996).

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* Catalysis Research Group, School of Chemistry and Physics, University of KwaZulu-Natal, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

955

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In most cases, the conventional Sonogashira coupling reactions are carried out under homogeneous conditions, which involve using palladium complexes such as [PdCl2(PPh3)2] or [Pd(PPh3)4] (3−5 mol%), a catalytic amount of a Cu salt (3-10 mol%) and an amine in large excess (Sonogashira, Tohda and Hagihara, 1975; Sonogashira, 2002; Rosa et al., 2015; Ciriminna et al., 2013). The use of copper salts as co-catalysts accelerates the Sonogashira reaction greatly; however, their presence in the reaction yields alkyne dimers (Sisodiya et al., 2015). This leads to the generation of an undesired homocoupling product of the alkyne upon oxidation. Thus, the reaction must be conducted under inert atmosphere if a Cu salt is to be used as a cocatalyst (Sisodiya et al., 2015). Furthermore, the above-mentioned standard conditions require a high loading of palladium and involve the use of phosphine ligands that are usually oxygen-sensitive (Cwik, Hell and Figueras, 2006). In addition, the efficient separation and subsequent recycling of homogenous palladium catalysts and ligand remains a challenge and an aspect of environmental and economic relevance (Roy, Senapati and Phukan, 2015). Intensive studies on heterogenization of Sonogashira coupling reactions with the objective of combining the benefits of both heterogeneous and homogeneous catalysis have been carried out during the last three decades (Abu-Reziq and Alper, 2012). In this direction, researchers have immobilised Pd on solid supports such as activated carbon (charcoal) (Rossy et al., 2014), zeolites

Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst (Djakovitch and Rollet, 2004), metal oxides (Cwik, Hell and Figueras, 2006; Hosseini-Sarvar, Razmi and Doroodmand, 2014; Roy, Senapati and Phukan, 2015; Kotadia et al., 2014) (MgO, ZnO, TiO2, ZrO2, Fe2O3, CeO2, .), clays (Borah and Dutta, (2013), alkaline earth salts (CaCO3, BaSO4, BaCO3, SrCO3) (Barros et al., 2008) and organic polymers (Ye and Yi, 2008; Kim et al., 2007). However, the leaching of palladium from the catalyst is still the main problem, even though most researchers have reported that their catalysts can be recycled by the redeposition of Pd onto the support (Redon, Peùa and Crescencio, 2014). This is due to the quasi-heterogeneous mechanism of the Sonogashira cross-coupling reaction via the solvated palladium clusters (Reay and Fairlamb, 2015). Hence, quasi-heterogeneous catalysis has developed into a new stream in catalysis and it is a promising approach for bridging heterogeneous and homogeneous catalysis (AbuReziq and Alper, 2012). Quasi-heterogeneous catalysis relies on the fact that nano-supports have a large surface areas, thus they can retain the activity and selectivity of the supported catalysts (Abu-Reziq and Alper, 2012; Wu et al., 2011). In this regard, we thought it would be worthwhile to develop a general catalytic system for a copper- and ligandfree quasi-heterogeneous Sonogashira coupling reaction catalysed by a nano Pd0¡02Ce0¡98O2- solid solution oxide. To the best of our knowledge, PdxCe1-xO2- based solid-solution oxides have not previously been investigated on the Sonogashira coupling reactions. Using an air, moisture and thermally stable quasi-heterogeneous Pd0.02Ce0.98O2- solid solution oxide should offer higher activity, simplicity of workup, recyclability and minimisation of metallic waste.

%B;59>; =?8 8@;:5;;@A? The structural and electronic properties of the Pd0¡02Ce0¡98O2- catalyst have been deduced earlier by X-ray diffraction (XRD), XPS, XANES, Raman spectroscopy, EXAFS and highresolution transmission electron microscopy (HR-TEM)

(Cwele et al., 2016).The characterisation results indicated that the prepared material was a solid solution oxide with a fluorite structure. The Rietveld refined XRD profile of the prepared catalyst and the XPS of the Pd(3d) are shown in Figure 1. The X-ray pattern corresponds to a monophasic Pd0.02Ce0.98O2- solid-solution oxide with a fluorite structure as reported earlier (Cwele et al., 2016). The insert in Figure 1 shows the XPS of the Pd(3d). The binding energy of Pd(3d5/2) and Pd(3d3/2) at 337.5 and 342.5 eV respectively, confirm that Pd is in the +2 oxidation state. The crystallite size and the lattice strain were estimated employing Williamson-Hall (W-H) plots and Equation [1]. The lattice strain in the Pd0.02Ce0.98O2- sample was four times greater than in the CeO2 sample (Table I). This indicates that lattice distortion occurs upon introduction of Pd2+ ions in the CeO2 lattice. The average crystallite size increased slightly with the incorporation of the Pd2+ ions into the CeO2 lattice. The average crystallite size of the blank CeO2 is 8 nm, while that of Pd0.02Ce0.98O2- increased to 13 nm. [1] The SEM and TEM images of CeO2 and Pd0.02Ce0.98O2- samples have very similar morphology (Figure 2). In both samples, the particles are present in the form of different sized lumps or flakes with rounded shaped structures. These types of structures are generated due to the nature of the solution combustion synthesis method, where gases escape giving rise to porosity. The BET surface area of the Pd0.02Ce0.98O2- sample was found to be 58 m2/g (Table I). The catalytic performance of the as-prepared Pd0.02Ce0.98O2- solid-solution precatalyst was evaluated in the Sonogashira coupling of phenylacetylene and various haloarenes and the results are listed in Table II. Optimisation of reaction conditions was conducted in order to develop a general catalytic system for the copperand ligand-free Pd0.02Ce0.98O2- catalysed Sonogashira

(@65<B ' 7B %@B>#B98 <B2@?B8 % 3=>>B<?; A2 $8+"+,)B+" ,- =?8 @>; $. A2 $80 8/ :A<B 9B#B9 <B6@A? 0@?;B>/

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Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst

(@65<B ,' @67> 1@:<A;:A34 0 /! .* 0 /! * 0)/ =?8 %- * 0 / =?=94;@; A2 >7B $8+"+,)B+" ,- :=>=94;>

Table I

$74;@:A:7B1@:=9 3<A3B<>@B; A2 >7B ;4?>7B;@;B8 1=>B<@=9;

CeO2 Pd0.02Ce0.98O2–δ

=>>@:B 3=<=1B>B< 0 /

=>>@:B ;><=@? 0 + /

)<4;>=99@>B ;@ B 0?1/

.5<2=:B =<B= 01, 6/

5.4221 5.4200

1.10 4.20

8 13

69 58

Table II

*22B:> A2 = &=;B A? >7B 1A8B9 .A?A6=;7@<= :<A;; :A539@?6 <B=:>@A? *?><4 1 2 3 4 5

=;B

%B=:>@A? >@1B 7

)A?#B<;@A?

Cs2CO3 Et3N K2CO3 NaOAc NaOH

24 1 24 24 24

62 100 66 21 48

coupling reaction. Solvent, temperature, alkyne loading, base and catalyst loading investigations were conducted using phenylacetylene (1) and iodobenzene (2) as model coupling partners (Scheme 1, Figure 3). We first conducted investigations to find a solvent system that is conducive for the Sonogashira coupling reaction (Figure 4). The solvent systems investigated were: acetonitrile (CH3CN), DMF, toluene (PhMe), ’solvent-less’ and H2O/CH3CN (1:3 v/v). Toluene gave the lowest conversion of 32% after 8 hours, while ’solvent-free’ and acetonitrile conditions gave relatively high conversion, of 68% and 80% respectively. DMF gave 100% conversion in 3 hours, while H2O/acetonitrile (1:3 v/v) went to completion in an hour. Hence, further investigations were conducted in H2O/acetonitrile (1:3 v/v). BÜhm and Herrmann (2000)

reported that the polarity and hydrogen bonding ability of the solvent are important in accelerating the reaction by stabilizing the ionic intermediates of the catalytic cycle. The solvent investigation results agree well with the above statement, since the reaction conducted in toluene has the longest reaction time and the lowest percentage conversion, while those conducted in aqueous acetonitrile had the shortest reaction time and gave 100% conversion of phenylacetylene. The BĂśhm and Herrmann findings also indirectly address the nature of catalysis for most Sonogashira coupling procedures (discussed later). The choice of base is considered crucial in Sonogashira coupling reactions. The base is used to neutralise the acid (HX) formed as a by-product in this reaction (BĂśhm and Herrmann, 2000). In addition, the base plays a role in inhibiting the generation of a homo-coupling product (Glaser-type reaction) and its strength plays a key role in the deprotonation of terminal alkynes (BĂśhm and Herrmann,

(@65<B '.:7B1B >7B 1A8B9 .A?A6=;7@<= :<A;;-:A539@?6 <B=:>@A? &B> BB? 37B?49=:B>49B?B 0 / =?8 @A8A&B? B?B 0,/ 9B=8@?6 >A &@37B?49=:B>49B?B 0 / VOLUME 117

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)=>=94;>

Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst

(@65<B '.A9#B?> @?#B;>@6=>@A? <B;59>; 5?8B< >7B 1A8B9 .A?A6=;7@<= :A539@?6 :A?8@>@A?;

2000; Komura, Nakamura and Sugi, 2008). Among the investigated bases, triethylamine was found to be the best base under our reaction conditions (Table II). The reaction went to completion in just an hour when triethylamine was used as base, but did not go to completion with the other bases investigated (K2CO3, NaOH, CH3COONa and Cs2CO3). It has been suggested that amines do not act as bases only, but also as weak coordinating ligands (Jutand, NÊgri and Principaud, 2005). Hence, beside the observed solubility problem when the inorganic bases were used, the fact that triethylamine can act as a coordinating base is thought to be a major reason for its superior performance. Three different temperatures were also investigated; 25, 50 and 82°C (Figure 5). It was found that the reactions were most rapid under refluxing conditions (82°C). Several experiments were conducted to find the lowest amount of the catalyst loading (based on Pd) that could efficiently catalyse the Sonogashira coupling reactions in

reasonable reaction times (Figure 6). The investigations were initiated with catalyst loading between 0.05-1 mol%. For loadings of 0.1-1 mol%, the reaction was found to be very fast and the reaction went to completion in an hour. This fast reaction makes it difficult to monitor the substitution and electronic effect caused by the aryl halides. Hence, a catalyst loading of 0.05 mol% was investigated; this then gave a slower reaction that went to completion in 3 hours. A combination of aryl halides and terminal alkynes was then studied to investigate the scope and limitations of our catalytic system under the obtained optimum reaction conditions. The investigations were initiated by reacting various iodoarenes with phenylacetylene, to investigate the electronic effect and the substitution position effect on the haloarenes. It was found that the aryl halides with electronwithdrawing groups in the para-position (Table III, entries 4, 5, 9 and 11) react more rapidly compared to those with

(@65<B '*22B:> A2 >B13B<=>5<B A? >7B 1A8B9 .A?A6=;7@<= :A539@?6 <B=:>@A?

(@65<B '*22B:> A2 :=>=94;> 9A=8@?6 A? >7B 1A8B9 .A?A6=;7@<= :A539@?6 <B=:>@A?

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Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst electron-donating groups in the same position (Table III, entries 2, 3 and 7). Hence, the electronic effect dictates the catalyst activity in the Sonogashira coupling reactions. It was also noted that the position of each functional group on the aryl halide also controls the activity of the catalyst. This is best demonstrated by the reactivity of 4-iodotoluene and 3-iodotoluene (Table III, entries 7 and 8); where the 3-iodotoluene reacts three times faster than 4-iodotoluene. These findings agree well with the reported influence of substituents in aryl halides on Sonogashira coupling reactions (Sonogashira, Tohda and Hagihara, 1975; Takahashi et al., 1980; Sonogashira, 2002). With the success obtained with aryl iodides, the reactions were extended to aryl bromides and aliphatic alkynes. However, bromoarenes, chloroarenes and aliphatic alkynes were found to be unreactive under the obtained optimum conditions. The stability of the C−Cl and C-Br bonds makes

chloroarenes and bromoarenes notoriously less reactive. Even under homogeneous copper-free coupling reaction conditions, aryl bromides give fairly low conversions (1060%) (Cwik, Hell and Figueras, 2006). Hence, the obtained optimum conditions are limited to iodoarenes and aromatic alkynes as coupling partners only. Performance comparisons of the present catalyst in Sonogashira cross coupling reactions of iodobenzene and phenylacetylene against other related Cu- and ligand-free catalytic systems are shown in Table IV.

B=:7@?6 =?8 <B:4:9=&@9@>4 >B;> A strong test for evaluating heterogeneity of a catalyst is the hot filtration test. To determine whether the reactions reported here were homogeneous or heterogeneous, a hot filtration test was performed for the Sonogashira coupling of

Table III

.A?A6=;7@<= :<A;;- :A539@?6 A2 37B?49=:B>49B?B @>7 #=<@A5; @A8A=<B?B; >A 6@#B ;5&;>@>5>B8 &@37B?49=:B>49B?B; 5?8B< A5< A3>@151 <B=:>@A? :A?8@>@A?;

A8A=<B?B

%B=:>@A? >@1B 07/

;A9=>B8 4@B98 0 /

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Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst Table IV

$B<2A<1=?:B :A13=<@;A?; A2 :=>=94;>; @? >7B <B=:>@A? A2 @A8A&B? B?B =?8 37B?49=:=>49B?B 5?8B< :A33B<- =?8 9@6=?8-2<BB :A?8@>@A?; *?><4 <B2

)=>=94;>

1(Roy, Senapati and Phukan, 2015) 2 (Komura, Nakamura and Sugi, 2008) 3 (Cwik, Hell and Figueras, 2006) 4 (Borah and Dutta, 2013) 5 (this paper)

%B=:>@A? :A?8@>@A?;

@1B 07/

@B98 0 /

Pd-CoFe2O4 (5 mol% Pd)

Ethanol, K2CO3, 70°C

6

90

Pd-2QC-MCM-41 (0.18 mol% Pd)

NMP, piperidine, 80°C

3

99

Pd-MgLa (1.5 mol% Pd)

DMF, Et3N, 90°C

10

90

Pd0/montmorill-onite (0.07 mol% Pd)

CH3CN, Et3N, 82°C

3

90

Pd0.02Ce0.98O2-δ (0.05 mol% Pd)

H2O/CH3CN (1:3), Et3N, 82°C

3

98

phenylacetylene and iodobenzene. The catalyst was filtered off from the reaction mixture after 5 minutes and the hot filtrate was then allowed to react further under similar conditions. The reaction was then monitored every 15 minutes for 2 hours. Assessment of the rate of iodobenzene conversion from the fresh reaction and the hot filtration test suggests that the activity in the fresh reaction can be attributed to the Pd species in the solution (Figure 7). The fresh catalyst and the hot filtrate have almost identical reaction profiles (Figure 7). An ICP analysis of the filtered solution confirmed the presence of about 0.5 ppm of Pd in the hot filtrate. Since the active catalytic phase was evidently the leached Pd, we attempted to determine the active Pd species using the mercury-poisoning test. Hg(0) is known to form amalgams with a variety of metals in their zero oxidation state (Reay and Fairlamb, 2015). Hence, if Pd dissociates from the solid Pd0.02Ce0.98O2- , it would bind with Hg(0) thereby quenching its catalytic activity. However, Pd in a raised oxidation state is not expected to be quenched by Hg(0) (Reay and Fairlamb, 2015). Thus, the mercury-poisoning test was used to establish whether bare Pd(0) did participate in catalysing the Sonogashira coupling reactions. To do the mercury poison test for the system described here, the Pd0.02Ce0.98O2- catalyst was filtered-off from the reaction mixture after a reaction time of 5 minutes and 2.5 mmol Hg(0) was added to the hot reaction mixture. The reaction mixture was then allowed to react further under similar conditions, while the activity was monitored every 15 minutes. It was found that no further reaction occurred after 15 minutes. Thus, the analysis confirms the presence of Pd(0) in the solution, which in turn gives insight into the reaction mechanism. Hence, the Pd0.02Ce0.98O2- -catalysed

Sonogashira coupling reaction is essentially a quasiheterogeneous catalytic process occurring over solvated Pd(0) clusters. The recycled catalyst showed about 60% conversion of iodobenzene in 2 hours. Hence, the recyclability test suggests that some of the leached Pd can redeposit on the CeO2 support at the end of the Sonogashira reaction and/or the Pd0.02Ce0.98O2- serves as an active catalyst reservoir.

$<A3A;B8 1B:7=?@;1 The exact mechanism by which the quasi-heterogeneous Pd0.02Ce0.98O2- -catalysed copper- and ligand-free Sonogashira coupling reaction occurs is still under investigation. However, the catalytic activity results, leaching and recyclability tests seem to indicate that the reaction follows a Pd(0)/Pd(2+) mechanism (Scheme 2 – Figure 8). We proposed that the first step in the reaction mechanism is ‘pre-activation’ which allows Pd2+0.02Ce0.98O2- to enter the catalytic cycle as Pd0/CeO2. The Pd(2+) ions in the fresh catalyst are reduced in situ by triethylamine and/or the solvent system (H2O/acetonitrile) to Pd(0). (Tougerti, Negri and Jutand, 2007; Jutand, NĂŠgri and Principaud, 2005) Since Pd(0) is more electron rich than Pd(2+), it allows for oxidative addition of the iodoarene, which marks the beginning of the catalytic cycle. The alkyne coordinates and forms a pi complex with Pd and then the desired Sonogashira coupling product is generated through the reductive elimination step.

)A?:95;@A? In conclusion, we have developed a general and simple process for Sonogashira coupling reactions using

(@65<B '%B=:>@A? 3<A2@9B; A2 2<B;7 :=>=94;>! 7A>-2@9><=>@A?! <B:4:9B8 :=>=94;> =?8 1B<:5<4-3A@;A? >B;>;

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Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst

(@65<B '.:7B1B , 3<A3A;B8 <B=:>@A? 1B:7=?@;1 2A< $8, +"+,)B+" ,- :=>=94;B8 5=;@-7A1A6B?BA5; .A?A6=;7@<= :<A;;-:A539@?6 <B=:>@A?;

*D3B<@1B?>=9 ;B:>@A? Cerium ammonium nitrate [(NH3)Ce(NO3)6, 99.9%], palladium chloride [PdCl2, 60%], urea [CH4N2O, 99.9%] were obtained from Sigma-Aldrich and were used without further purification.

The monophasic Pd0.02Ce0.98O2- solid-solution oxide was prepared using a one-step urea-assisted solution combustion synthesis method described earlier (Cwele et al., 2016). The catalyst synthesis method involved preparation of a redox combustion mixture composed of stoichiometric amounts of metal precursors [(NH4)2Ce(NO3)6 and PdCl2] and urea [NH2CONH2] in the ratio of 1.0:3.77 respectively, dissolved in water. The solution was then stirred at 150°C for 10 minutes to evaporate water and reduce its volume to approximately 20 mL. The boiling solution was then introduced into a muffle furnace pre-heated at 400°C and was kept in the furnace overnight. A light brown solid was obtained.

A Bruker D8 Advance diffractometer, equipped with a XRK900 in-situ cell and a Cu K source ( = 1.5406 Ă…) was used to record the powder X-ray diffraction patterns of the samples. The structures were refined by the Rietveld method using the Full Prof Suite-2000 program. The average crystallite size (D) and lattice strain ( ) of CeO2 and Pd0.02Ce0.98O2- were estimated from the modified Rietveld method and Williamson-Hall (W-H) plots.

ICP-OES was performed using a Perkin Elmer optical emission spectrometer Optima 5300 DV. Standards (1000 ppm Ce and Pd) were purchased from Fluka. Brunauer−Emmett−Teller (BET) surface area measurements were determined using a MicroMetrics TriStar 3000 porosimeter with N2 as probe gas. About 0.4 g of each powder sample was degassed overnight at 200°C using a Micromeritics FlowPep 060 instrument prior to analysis. SEM images were obtained with a Jeol JSM-6100 scanning electron microscope using a Bruker signal processing unit detector. The analysis was performed at random points along the surface of the catalyst. The samples were first mounted on aluminium stubs using double-sided carbon tape; they were then coated with gold using a Polaron E5100 coating unit. For TEM analysis, the samples were viewed on a Joel JEM-1010 electron microscope. For high-resolution TEM (HR-TEM) and scanning electron microscopy (STEM) analysis, the samples were viewed on a Joel JEM-2100 electron microscope and the images captured were analysed using iTEM software. The powder samples were ultrasonically dispersed in ethanol and supported on a perforated carbon film mounted on a copper grid prior to analysis.

A dry two-necked pear-shaped flask containing a stirrer bar, a condenser, 2 mL of H2O and 6 mL of acetonitrile was charged with aryl halide (2.3 mmol), phenylacetylene (2 eq.), triethylamine (5 eq.) and catalyst Pd0.02Ce0.98O2- (0.05 mol% Pd). The reaction mixture was stirred and (usually) heated to 82°C and its progress was monitored by GC and GC-MS. The iodoarene conversion was used to estimate the catalytic activity, using biphenyl as an internal standard. After the reaction had gone to completion, the reaction mixture was filtered and the filtrate was extracted with ethyl acetate and brine. The organic layer was then evaporated under reduced pressure and the residue was purified by flash chromatography on silica gel using VOLUME 117

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Pd0.02Ce0.98O2- as a precatalyst. The catalyst was found to be very effective in a water/acetonitrile solvent system at 82°C under ligand- and Cu-free reaction conditions. A wide range of iodorenes were coupled to phenylacetylene using this catalytic route. It was also found that the Pd0.02Ce0.98O2- catalysed Sonogashira cross-coupling reactions are essentially a quasi-heterogeneous catalytic process occurring over solvated Pd(0) clusters. As a result, the recovered catalyst loses some activity upon recycle.

Pd0.02Ce0.98O2- : a copper- and ligand-free quasi-heterogeneous catalyst EtOAc/hexane (8/2) as an eluent. The structure of the coupling product was confirmed by 1H and 13C NMR spectroscopy and the results were consistent with those reported in the literature for substituted biphenylacetylenes (Djakovitch and Rollet, 2004).

The catalyst used in the first run was separated by centrifugation, washed with 5 mL acetonitrile, dried at 60°C and reused as described for the fresh catalyst.

:C?A 9B86B1B?>; We would like to express our gratitude to Mintek and the Department of Science and Technology (Advanced Metals Initiative program) for financial support. The authors would also like to thank Dr A.S. Mohamed and Dr S. Singh for helpful discussions.

KING, A.O. and YASUDA, N. 2004. Palladium-catalyzed cross-coupling reactions in the synthesis of pharmaceuticals. Topics in Organometallic Chemistry, vol. 6. pp. 205–245. KOMURA, K., NAKAMURA, H. and SUGI, Y. 2008. Heterogeneous copper-free Sonogashira coupling reaction of terminal alkynes with aryl halides over a quinoline-2-carboimine palladium complex immobilized on MCM-41. Journal of Molecular Catalysis A: Chemical, vol. 293, nos. 1–2. pp. 72–78. KOTADIA, D.A., PATEL, U.H., GANDHI, S. and SONI, S. S. 2014. Pd doped SiO2 nanoparticles: an efficient recyclable catalyst for Suzuki, Heck and Sonogashira reactions. RSC Advances, vol. 4, no. 62. pp. 32826–32833. MUJAHIDIN, D. and DOYE, S. 2005. Enantioselective Synthesis of (+)-(S)Laudanosine and (–)-(S)-Xylopinine. European Journal of Organic Chemistry, vol. 13. pp. 2693–2693. REAY, A.J. anD FAIRLAMB, I.J.S. 2015. Catalytic C–H bond functionalisation chemistry: the case for quasi-heterogeneous catalysis. Chemical Communications, vol. 51, no. 91. pp. 16289–16307. REDON, R., PEÑA, N.G. and CRESCENCIO, F.R. 2014. Leaching in metal nanoparticle catalysis. Recent Patents on Nanotechnology, vol. 8, no. 1. pp. 31–51.

%B2B<B?:B; ABU-REZIQ, R. and ALPER, H. 2012. Magnetically separable base catalysts: heterogeneous catalysis vs. quasi-homogeneous catalysis. Applied Sciences, vol. 2, no. 2. pp. 260–276. BARROS, J.C., DE SOUZA, A.L.F., DE LIMA, P.G., DA SILVA, J.F.M. and ANTUNES, O.A.C. 2008. Selectivity studies in the reaction between iodobenzene and phenylacetylene: Sonogashira coupling vs. hydroarylation. Applied Organometallic Chemistry, vol. 22, no. 5. pp. 249–252. BEUTLER, U., MAZACEK, J., PENN, G., SCHENKEL, B. and WASMUTH, D. 1996. Development of a new environmentally friendly production process for terbinafine. Chimia, vol. 50, no. 4. pp. 154–156. BĂ–HM, V.P.W. and HERRMANN, W.A. 2000. A copper-free procedure for the palladium-catalyzed Sonogashira reaction of aryl bromides with terminal alkynes at room temperature. European Journal of Organic Chemistry, vol. 22. pp. 3679-3681. BORAH, B.J. and DUTTA, D.K. 2013. In situ stabilization of Pd0-nanoparticles into the nanopores of modified montmorillonite: Efficient heterogeneous catalysts for Heck and Sonogashira coupling reactions. Journal of Molecular Catalysis A: Chemical, vol. 366. pp. 202–209. CIRIMINNA, R., PANDARUS, V., GINGRAS, G., BÉLAND, F., DEMMA CARĂ€, P. and PAGLIARO, M. 2013. Heterogeneous Sonogashira coupling over nanostructured siliacat Pd(0). ACS Sustainable Chemistry & Engineering, vol. 1, no. 1. pp. 57–61. CWELE, T., MAHADEVAIAH, N., SINGH, S., FRIEDRICH, H.B., YADAV, A.K., JHA, S.N., BHATTACHARYYA, D. and SAHOO, N.K. 2016. CO oxidation activity enhancement of Ce0.95Cu0.05O2− induced by Pd co-substitution. Catalysis Science & Technology, vol. 6, no. 22. pp. 8104–8116. CWIK, A., HELL, Z. and FIGUERAS, F. 2006. A copper-free Sonogashira reaction using a Pd/MgLa mixed oxide. Tetrahedron Letters, vol. 47, no.18. pp. 3023–3026. DJAKOVITCH, L. and ROLLET, P. 2004. Sonogashira cross-coupling reactions catalysed by copper-free palladium zeolites. Advanced Synthesis & Catalysis, vol. 346, nos. 13–15. pp. 1782-1792. HOSSEINI-SARVARI, M., RAZMI, Z. and DOROODMAND, M.M. 2014. Palladium supported on zinc oxide nanoparticles: Synthesis, characterization and application as heterogeneous catalyst for Mizoroki–Heck and Sonogashira reactions under ligand-free and air atmosphere conditions. Applied Catalysis A: General, vol. 475. pp. 477–486. JUTAND, A., NÉGRI, S. and PRINCIPAUD, A. 2005. Formation of ArPdXL(amine) complexes by substitution of one phosphane ligand by an amine in transArPdX(PPh3)2 complexes. European Journal of Inorganic Chemistry, vol. 4. pp. 631–635. KIM, J.H., LEE, D.H., JUN, B.H. and LEE, Y.S. 2007. Copper-free Sonogashira

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ROSA, D.S., ANTELO, F., LOPES, T.J., DE MOURA, N.F. and ROSA, G.R. 2015. Effects of solvent, base and temperature in the optimisation of a new catalytic system for sonogashira cross-coupling using NCP pincer palladacycle. Quimica Nova, vol. 38, no. 5. pp. 605–608. ROSSY, C., MAJIMEL, J., DELAPIERRE, M.T., FOUQUET, E. and FELPIN, F.X. (2014). On the peculiar recycling properties of charcoal-supported palladium oxide nanoparticles in Sonogashira reactions. Applied Catalysis A: General, vol. 482. pp. 157–162. ROY, S., SENAPATI, K.K. and PHUKAN, P. 2015. Direct use of nanoparticles as a heterogeneous catalyst: Pd0-doped CoFe2O4 magnetic nanoparticles for Sonogashira coupling reaction. Research on Chemical Intermediates, vol. 41, no. 8. pp. 5753–5767. SISODIYA, S., WALLENBERG, L.R., LEWIN, E. and WENDT, O.F. 2015. Sonogashira coupling reaction over supported gold nanoparticles: Influence of support and catalyst synthesis route. Applied Catalysis A: General, vol. 503, pp. 69–76. SONOGASHIRA, K. 200. Development of Pd–Cu catalyzed cross-coupling of terminal acetylenes with sp2-carbon halides. Journal of Organometallic Chemistry, vol. 653, no. 2. pp. 46–49. SONOGASHIRA, K., TOHDA, Y. and HAGIHARA, N. 1975. A convenient synthesis of acetylenes: catalytic substitutions of acetylenic hydrogen with bromoalkenes, iodoarenes and bromopyridines. Tetrahedron Letters, vol. 16, no. 50. pp. 4467–4470. TAKAHASHI, S., KUROYAMA, Y., SONOGASHIRA, K. and HAGIHARA, N. 1980. A convenient synthesis of ethynylarenes and diethynylarenes. Synthesis, vol. 8. pp. 627–630. TORBORG, A. and BELLER, M. 2009. Recent applications of palladium-catalyzed coupling reactions in the pharmaceutical, agrochemical and fine chemical industries. Advanced Synthesis & Catalysis, vol. 351, no. 18. pp. 3027–3043. TOUGERTI, A., NEGRI, S. and JUTAND, A. 2007. Mechanism of the copper-free palladium-catalyzed Sonagashira reactions: Multiple role of amines. Chemistry-A European Journal, vol. 13, no. 2. pp. 666–676. WU, Y., WANG, D., ZHAO, P., NIU, Z., PENG, Q. and LI, Y. (2011). Monodispersed Pd−Ni nanoparticles: composition control synthesis and catalytic properties in the Miyaura−Suzuki. Inorganic Chemistry, vol. 50, no. 6. pp. 2046–2048. YE, Z.W. and YI, W.B. (2008). Polymer-supported palladium perfluorooctanesulfonate [Pd(OPf)2]: A recyclable and ligand-free palladium catalyst for copper-free Sonogashira coupling reaction in water under aerobic conditions. Journal of Fluorine Chemistry, vol. 129, no. 11. pp. 1124–1128.

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a6

First-principles studies of Fe-Al-X (X=Pt, Ru) alloys by C.S. Mkhonto, H.R. Chauke and P.E. Ngoepe

The Fe-Al based alloys have recently attracted a lot of attention due to their excellent resistance to oxidation at high temperatures. However, they suffer limited room temperature ductility and a sharp drop in strength above 600°C. The current study employed a density functional theory approach to investigate the stability of FeAl-X alloys. We employed virtual crystal approximation to model various atomic concentrations (0 X 5) of both Pt and Ru; this will allow more precise predictions on the materials’ behaviour. Density of states was used to describe the behaviour of each phase near the Fermi level; these phases were observed at different percentage compositions. The FeAl composition is most favourable since it displays positive shear moduli, condition of mechanical stability. Addition of Pt and Ru was found to significantly improve the ductility of the Fe-AlX compound for 0.2 and 0.5 at.% compositions, respectively. C@* :</6 FeAl-X alloys, DFT, heats of formation, thermodynamic stability, density of states, elastic constants, X-ray diffraction patterns.

B;><:/3&>9:; The intermetallic iron-aluminium system has attracted a large amount of research since it possesses good mechanical properties, low density and low cost, as well as easy access to the raw materials (Couperthwaite, Cornish and Mwamba, 2016). It was also reported that the Fe-Al alloys are promising materials, due to their good refractoriness, oxidation and corrosion resistance and good ductility at room temperature (Couperthwaite, Cornish and Mwamba, 2016). However, these materials suffer limited room temperature ductility and a sharp drop in strength above 600°C, which makes them less suitable for use as structural materials (Li et al., 2016). Previous experimental work investigated the effect of precious metals (Pd, Ag, Ru, Pt) additions on the structure, oxidation and corrosion properties of a Fe-40 at.% Al (Fe-Al) alloy. It was reported that additions of more than 0.5 at.% precious metal did not improve the oxidation and corrosion properties of the materials and in some cases even decreased the resistance to corrosion (Couperthwaite Cornish and Mwamba, 2016; Li et al., 2016).

VOLUME 117

* Materials Modelling Centre, School of Physical and Mineral Sciences, University of Limpopo, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

963

*;:)696

Four alloys were produced through mechanical alloying i.e. FeAl-0.2 at.% Pd, FeAl-0.2 at.% Ru, FeAl-0.5 at.% Ag, FeAl-0.5 at.% Pt. It was also found that the additions of Ru and Pt were crucial to the oxidation and corrosion properties, with Ru being considered the most favourable. In previous work, it was established that additions of 0.2 at.% Ru to a Fe-40 at.% Al alloy improved the corrosion and oxidation resistance. Furthermore, the findings revealed that the non-equilibrium processing significantly refined the grain size of the material. The sintered material had a higher hardness than the as-cast material and the change in grain size did not significantly affect the oxidation and corrosion resistance (Chou et al., 2006). Other research focused on the mechanically alloyed powder coated onto a mild steel substrate (approx. 5–10 m thick) at a low gas pressure of 10 bar and temperature of 500°C. It was found that the coated materials effectively enhanced the oxidation and corrosion resistance (Chou et al., 2006). Iron aluminides based on Fe3Al and FeAl are highly oxidation and corrosion resistant and have potential for elevated temperature structural applications (Kant et al., 2016). Carbon is an important alloying element in Fe3Al as it increases strength and creep resistance, as well as resistance to environmental embrittlement (Colinet, 2003). However, addition of carbon to FeAl has not been successful as it leads to precipitation of graphite which causes decrease in strength (Colinet, 2003).

First-principles studies of Fe-Al-X (X=Pt, Ru) alloys In this study, DFT was used to investigate the structural, electronic and mechanical properties of Ru and Pt doped FeAl intermetallic. We employed the virtual crystal approximation (VCA) which allowed calculations on disordered systems to be carried out at the same cost as calculations for ordered structures (Ramer and Rappe, 2000). Small amounts of up to 0.5 at.% Pt and up to 0.2 at.% Ru were investigated to complement the previous experimental findings) (Couperthwaite, Cornish and Mwamba, 2016). The study modelled various atomic concentrations (0 X 5) of both Pt and Ru; to allow more precise predictions and understanding of the electronic and elastic behaviour of the FeAl compounds. The model is represented in Figure 1, showing possible doping of metal (M=Pt, Ru) atoms on either Al or Fe sublattices (Marker et al., 2013; Inden and Pepperhoff, 1990).

doped with either Ru or Pt atoms. The metal concentration of 0.2 at.% Ru and 0.5 at.% Pt was considered (Couperthwaite, Cornish and Mwamba, 2016) to determine the structural and thermodynamic stability and to understand the electronic and elastic property signatures. Firstly, the most energetically favourable structure was determined by checking various metal doping concentrations on different sublattices as shown in Table I. The lattice parameter or volume mismatches for Pt and Ru were insignificant, since all the composition gave lattice parameters very close to one another (ranging between 0.280 nm and 0.290 nm). We have predicted that the Pt and Ru prefer the Fe sublattice with a composition of Fe49.8Al50Ru0.2 and Fe49.5Al50 Pt0.5. This analysis is deduced from the heats of formation of –0.455 eV and –0.431 eV per atom for Ru and Pt doping, respectively.

:-)3>?>9:;?=A-@>":/:=:(*

In order to evaluate the elastic stability of the ternary systems, we calculated the elastic constants (Cij) for Fe50Al49.8Ru0.2, Fe49.8Al50Ru0.2, Fe50Al49.5Pt0.5, Fe49.5Al50Pt0.5, Fe49.9Al49.9Ru0.2 (50:50) and Fe49.75Al49.75Pt0.5 (50:50) dopant at different concentrations as shown in Table I. Note that the symmetry was unchanged and only three independent elastic constants (C11, C12 and C44) have been found for the cubic lattice. The mechanical stability criteria of cubic system (Mehl, Klein and Papaconstantopoulos, 1994) is given as:

Density functional theory (DFT) (Kohn and Sham, 1965) within generalised gradient approximation (GGA-PBE) exchange-correlation functional (Perdew, Burke and Ernzerhof, 1996) was used to study the Fe-Al alloys. We employed the plane wave pseudopotential method as implemented in CASTEP cod (Milman et al., 1999; Hedin and Lundqvist, 1971). An energy cutoff of 500 eV was used, as it was sufficient to converge the total energy of the B2 FeAl phase. The Brillouin zone integrations were performed for suitably large sets of k-points. We used a 10 Ă— 10 Ă— 10 k-points before and after doping. In the calculation of elastic constants and density of states, a k-spacing of 0.2 was used. Optimisation of structural parameters (atomic positions and lattice parameters) was achieved by minimisation of forces and stress tensors. Initially, the optimised binary B2 FeAl structure gave equilibrium lattice parameter of 0.2852 nm, in good agreement with the experimental value of 0.2908 nm (Breuer et al., 2001). The predicted heat of formation was 0.333 eV per atom compared to the experimental value of 0.376 eV per atom (Breuer et al., 2001).

and shear is

$@63=>6A?;/A/96&3669:; Figure 1 shows the binary B2 FeAl structure and VCA model

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Table I

39=9 <93-A=?>>9&@A)?<?-@>@<6 A :=3-@A?;/A"@?>6A:'A':<-?>9:;A8 7A:'A>"@A4@ 0= A?==:*6 :-):69>9:;

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,>

?>>9&@A)?<?-@>@<A8;-7

:=3-@A8;- 7

8@ A)@<A?>:-7

-0.453

Fe49.90Al49.90Ru0.2

0.2

-

0.2848

2.3107

Fe49.80Al50Ru0.2

0.2

-

0.2851

2.3185

-0.455

Fe50Al49.80Ru0.2

0.2

-

0.2848

2.3108

-0.452

Fe49.75Al49.75Pt0.5

-

0.5

0.2852

2.3220

-0.429

Fe49.50Al50Pt0.5

-

0.5

0.2852

2.3197

-0.431

Fe50Al49.50Pt0.5

-

0.5

0.2853

2.3220

-0.427

964

VOLUME 117

First-principles studies of Fe-Al-X (X=Pt, Ru) alloys All independent constants were positive and satisfied the stability conditions of the cubic lattice. Furthermore, the calculated shear moduli (C’) is positive, indicating the all compounds under considerations were elastically stable (Coudert and Mouhat, 2014). The elastic stability criteria also led to a restriction on the magnitude of B. Since B is a weighed average of C11 and C12 and stability requires that C12 be smaller than C11, we are then left with the result that B is required to be intermediate in value between C11 and C12: C12 <B<C11. We predicted high values of B for all systems with the lowest values corresponding to stable compounds. Thus, from the predicted B values, we determined the ductility and brittleness of these compounds from the ratio of bulk to shear moduli. Pugh (1954) proposed the B/G ratio predicted the ductility (> 1.75) or brittleness (< 1.75). For cubic Fe-Al alloys in Table II, we observed that B/G < 1.75, suggesting brittleness of the material. Note also that the negative values

of B/G also reflect instability of the corresponding compounds, which is not the case for the Fe-Al-X alloys. We noticed that the ductility is slightly enhanced for Fe49.50Al50Pt0.5 and Fe49.80Al50Ru0.2 systems (Breuer et al., 2001).

The electronic density of states (DOS) was calculated to mimic the stabilities by observing the behaviour of electronic states near the Fermi level. This approach has been used effectively in many studies and is mainly used to confirm or correlate the thermodynamic stability of intermetallic compounds (Mahlangu et al., 2013; Phasha et al., 2010). Firstly, we showed the total and partial density of states for the binary FeAl system in Figure 2a. We observed that the structure was characterised by a pseudogap near the Fermi level (confirming a metallic behaviour character). A more

Table II

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## 8!,?7

#% 8!,?7

.. 8!,?7

A8!,?7 # %A8 ## #%7

8!,?7

8!,?7

8!,?7A

Fe49.90Al49.90Ru0.2

283.0

141.4

148.3

212.3

110.219

187.892

1.705

Fe49.80Al50Ru0.2

279.4

140.6

147.4

209.1

108.942

186.885

1.715

Fe50Al49.80Ru0.2

282.2

140.7

148.2

211.8

110.144

188.590

1.712

Fe49.75Al49.75Pt0.5

282.5

147.9

147.9

208.5

107.828

187.304

1.737

Fe49.50Al50Pt0.5

278.9

139.9

147.5

208.9

101.799

186.243

1.707

Fe50Al49.50Pt0.5

282.0

139.3

147.9

212.4

103.487

186.917

1.693

VOLUME 117

965

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First-principles studies of Fe-Al-X (X=Pt, Ru) alloys noticeable Fe d-peak was observed, which forms a strong hybridisation with the Al p-orbital. More importantly, we see that the Fermi level fell slightly on the left of the pseudogap, which signifies electronic stability in agreement with the predicted heats of formation (i.e. good correlation).

Figures 2b and 2c show the total and partial density of states for Fe49.5Al50Pt0.5 and Fe49.8Al50Ru0.2, respectively. It is clear from the plot that doping with Pt and Ru slightly changed the behaviour of the electronic structure, in particular, shifting of the Fermi level with respect to the pseudogap. Note that the two plots show 0.5 at.% Pt and 0.2 at.% Ru doped on the Fe sublattice, noticeably the partial DOS for Al, were similar (row 2 of Figures 2b and 2c), while those for Fe/Pt (in Figure 2b) and for Fe/Ru (in Figure 2c) were different. A small peak at about –0.5 eV was sharper for the Ru-doped system than for Pt. More importantly, we noticed that the Fermi level slightly fell in the pseudogap (at E–EF = 0) for both 0.5 at.% Pt and 0.2 at.% Ru, similar to the binary phase (Figure 2a).

Accordingly, doping with Pt on the Fe sublattice would be a preference for the given concentration, since it is more stable. The total and partial DOS for the doped structures are compared in Figure 3. We observed similar trends as those in Figure 2, the plots only shows significant difference with the appearance of peaks at about –-0.5 eV, which was attributed to either Pt or Ru additions.

We observed the X-ray diffraction patterns for binary FeAl and the stable compounds (Fe49.5Pt0.5 and Fe49.8Al50Ru0.2), as shown in Figure 4. These plots displayed similar peaks before and after doping. However, the intensity of doped Ru system was very similar to the binary FeAl phase. Most notable was the high intensity peak at about 32 a.u and 70 a.u. This XRD pattern confirmed the structure and the fact that it did not change was attributed to the small amounts of Pt and Ru that dissolved into the structure.

3--?<*A?;/A&:;&=369:; The equilibrium lattice parameter, heats of formation, elastic

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First-principles studies of Fe-Al-X (X=Pt, Ru) alloys Research Chair Initiative of the Department of Science and Technology and the National Research Foundation is highly appreciated.

$@'@<@;&@6 ATABAKI, M.M., NIKODINOVSKI, M., CHENIER, P., MA, J., HAROONI, M. and KOVACEVIC R. (2014). Welding of aluminium to steels: an overview. Journal for Manufacturing Science and Production, vol. 14. pp. 2191–4184. BREUER, J., GRÃœN, A., SOMMER, F. and MITTEMEIJER, E.J. 2001. Enthalpy of formation of B2-Fe1−xAlx and B2-(Ni,Fe)1−xAlx. Metallurgical and Materials Transactions B, vol. 32. pp. 913–918. CHOU, S., HUANG, J., LII, D. and LU, H. 2006. The mechanical properties of Al2O3/aluminum alloy A356 composite manufactured by squeeze casting. Journal of Alloys and Compounds, vol. 419. pp. 98–102. COLINET, C. 2003. ab-initio calculation of enthalpies of formation of intermetallic. Intermetallics, vol. 11. pp. 1095–1102. COUPERTHWAITE, R.A., CORNISH, L.A. and MWAMBA, I.A. 2016. Cold-spray coating of Fe-40 at. % Al alloy with additions of ruthenium. Journal of the Southern African Institute of Mining and Metallurgy,vol. 116. pp. 927–934. HEDIN, L. and LUNDQVIST, B.I. 1971. Explicit local exchange-correlation potentials. Journal of Physics C: Solid State Physics, vol. 4. pp. 2064–2082. INDEN, G. and PEPPERHOFF, W. 1990. Experimental study of the order: disorder transition in bcc Fe-Al alloys. Zeitschrift für Metallkunde, vol. 81. pp. 770–773. KANT, R., PRAKASH, U., AGARWALA, V. and STYA PRASAD, V.V. 2016. Effect of carbon and titanium additions on mechanical properties of B2 FeAl. Transactions of the Indian Institute of Metals, vol. 69. pp. 845–850. KOHN, W. and SHAM, L.J. 1965. Self-consistent equations including exchange and correlation effects. Physical Review, vol. 140. pp. 1133–1138. LI, X., SCHERF, A., HEILMAIER, M. and STEIN, F. 2016. The Al-rich part of the FeAl phase diagram. Journal of Phase Equilibria and Diffusion, vol. 39. pp. 162–173. MAHLANGU, R., PHASHA, M.J., CHAUKE, H.R. and NGOEPE, P.E. 2013. Structural, elastic and electronic properties of equiatomic PtTi as potential hightemperature shape memory Alloy. Intermetallics, vol. 33. pp. 27–32.

properties and electronic structure of the B2 FeAl phase were determined using ab initio calculations. Interestingly, the B2 FeAl at 50:50 was found to be energetically and mechanically stable over the other phases at different compositions. These phases exist in the range between 23 and 55 atomic percent Al of the experimental phase diagram (Atabaki et al., 2014). The shear modulus (C’) of the B2 FeAl phase was found to be positive, fulfilling the condition of stability. The DFT results were in good agreement with the experimental findings (Couperthwaite, Cornish and Mwamba, 2016). It was found that the FeAl structure was more stable (lowest heats of formation) and VCA showed preference of doping on Fe rather than Al sublattices: Fe48.8Al50Ru0.2 and Fe49.5Al50Pt0.5; and finally the DOS showed that the Fermi level fell in the pseudogap which demonstrates the condition of stability.

0& ;: =@/(@-@;> The work was carried at the Materials Modelling Centre, University of Limpopo. The support of the South African

MEHL, M.J., KLEIN, B.M. and PAPACONSTANTOPOULOS, D.A. (1994). Principles. Intermetallic Compounds, vol. 1. Westbrook J.H. and Fleischer R.L. (eds.). Wiley. MILMAN, V., WINKLER, B., WHITE, J.A, PICKARD, C.J., PAYNE, M.C., AKHMATSKAYA, E.V. and NOBES, R.H. 1999. Electronic structure, properties, and phase stability of inorganic crystals: a pseudopotential plane-wave study. International Journal of Quantum Chemistry, vol. 77. pp. 895–910. MOUHAT, F. and COUDERT, F. 2014. Neccesary and sufficient elastic conditions in various crystal systems. Physical Review B, vol. 90. pp. 224104–224107. PERDEW, J.P., BURKE, K. and ERNZERHOF, M. 1996. Generalized gradient approximation made simple. Physical Review Letters, vol. 77. pp. 3865–3868. PHASHA, M.J., NGOEPE, P.E., CHAUKE, H.R., PETTIFOR, D.G. and NGUYEN-MANN, D. 2010. Link between structural and mechanical stability of fcc- and bccbased ordered Mg-Li alloys. Intermetallics, vol. 18. pp. 2083–2089. PUGH, S.F. 1954. XCII. Relations between the elastic moduli and the plastic properties of polycrystalline pure metals. The London, Edinburgh, Dublin Philosophical Magazine and Journal of Science, vol. 45, no. 367. pp. 823–843. RAMER, N.J. and RAPPE, A.M. 2000. Application of a new virtual crystal approach for the study of disordered perovskites. Journal of Physics and Chemistry of Solids, vol 61. pp. 315–320. VOLUME 117

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49(3<@A. $ A)?>>@<;6A:'A 9;?<*A4@0=A?;/A/:)@/A4@.+510=12,>251A?;/ 4@.+5 0=12$325%

MARKER, M.C.J., DUARTE, L.I., LEINENBACH, C. and RICHTER, K.W. 2013. Characterization of the Fe-rich corner of Al-Fe-Si-Ti. Intermetallics, vol. 39. pp. 38–49.

The Society of Mining Professors (SOMP) in collaboration with the Mining Engineering Education South Africa (MEESA) and The South African Institute of Mining and Metallurgy (SAIMM) is proud to host

Society of Mining Professors 6th Regional Conference 2018 Overcoming challenges in the Mining Industry through sustainable mining practices

12–13 March 2018 — Conference 14 March 2018— Technical Visit Birchwood Hotel and Conference Centre, Johannesburg, South Africa

The Mining Engineering Education South Africa (MEESA) will host the Society of Mining Professors (SOMP) 6th Regional Conference with the theme: Overcoming challenges in the Mining Industry through sustainable mining practices. The Society of Mining Professors is a vibrant Society representing the global academic community and committed to making a significant contribution to the future of the minerals disciplines. The main goal of the Society is to guarantee the scientific, technical, academic and professional knowledge required to ensure a sustainable supply of minerals for mankind. The Society facilitates information exchange, research and teaching partnerships and other collaborative activities among its members. MEESA is comprised of the School of Mining Engineering at the University of Witwatersrand, the Department of Mining Engineering at the University of Pretoria, the Department of Mining Engineering at the University of Johannesburg and the Department of Mining Engineering at the University of South Africa. The 6th Regional Conference gives a platform to academics, researchers, government officials, Minerals Industry professionals and other stakeholders an opportunity to interact, exchange and analyse the challenges and opportunities within the Minerals Industry. For any country to develop technologically and economically there must be a strong link between its industry, government & academic institutions. This conference will put together the role-players of the Mineral Industry from within and outside South Africa. The main aim of this conference is to facilitate information exchange. It is known that the mining industry is currently faced with big challenges ranging from the technical skills shortage, deep ore bodies, declining ore grades, challenges linked to processing ores with complex mineralogy, water quality and supply to the ever escalating energy costs and sustainability amongst others (Musiyarira et al., 2014). To address some of these challenges there must be a strong link between its industry, government & academic institutions. This will only happen when all the role-players collaboratively work together. The set-up of this conference is such that it allows the interaction between the Minerals Industry players and the academics.

The conference will feature peer reviewed technical presentations from academic, government and Industry professionals on a wide range of topics. While outlining the Conference programme, great emphasis is laid on participants’ interaction in addition to the presentations. The conference will be structured as follows: N N N N N

Two-day technical programme with peer-reviewed papers Relevant discussion workshops Field trips and site tours Networking opportunities Keynote lecturers

The Conference is being organised by Society of Mining Professors (SOMP) in collaboration with the Mining Engineering Education South Africa (MEESA) and The Southern African Institute of Mining and Metallurgy (SAIMM). The Conference presenters are well-known and highly respected experts in their fields, and will cover a wide range of topics. Presentations will be followed by discussions. The Conference will be of benefit to the Minerals Industry professionals, academics, non-government, government officers and other stakeholders. For all enquiries please contact: Chair: Associate Prof Rudrajit Mitra Telephone: +27 (011) 717 7572 | Email: rudrajit.mitra@wits.ac.za Head of SOMP Committee on Capacity Building: Dr Harmony Musiyarira Head of Conferencing: Camielah Jardine SAIMM, P O Box 61127, Marshalltown 2107 Tel: +27 (0) 11 834-1273/7 ¡ E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za

S ociety of Mining Professors Societät der Bergbaukunde

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a7

PGMs: A cornucopia of possible applications by L.A. Cornish

South Africa has been mining PGMs for many years and has a vested interest in the usefulness of these metals, although originally, there was little thought about downstream beneficiation and some of the products were later purchased, in effect repurchasing the raw materials. However, for about 20 years, it has been realised that in order to build wealth sustainably, at least some of the beneficiation and development has to be done in South Africa. This is because for each stage of beneficiation, the value of the commodity increases and so to buy the PGMs back in a later state is uneconomical. Also, the industries need to be established here in South Africa, so that South Africans can work in them and become part of the workforce and generate their own wealth and sustainability for their families. This is part of the rationale of the AMI. In order to drive this, as well as a viable commercial strategy, with viable and wanted products, there has to be something different in the production mixture, for example, the product has to be better, new or cheaper. This is where research is used to try and develop new capabilities, or to improve them. PGMs, application, product development.

South Africa has been mining PGMs for many years and has been the leading supplier of PGMs because of the richness of the Bushveld Complex. Originally, there was little thought about downstream beneficiation, with more emphasis being placed on exporting the metals. Some of the products were later purchased, which is uneconomical, because for each stage of beneficiation, the value of the commodity increases. However, for about 20 years, it has been realised that in order to build wealth sustainably, South Africa must have a vested interest in the usefulness of these metals and that at least some of the beneficiation and development has to be done in locally. Also, the industries need to be established here in South African order to create emploment oportunities, so that South Africans can generate their own wealth and sustainability for their families. This is part of the rationale of the Advanced Metals Initiative (AMI). In order to drive this, as well as a viable commercial strategy, with viable and wanted products, there has to be something different

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Platinum-based alloys possess high melting points, good thermal stability and thermal shock resistance and good corrosion and oxidation resistance and different approaches to strengthening. The alloys were established by the 1980s (Hammer and Kaufmann, 1982; Heywood, 1988; Heraeus, 2011). There are several applications where the high electrical

* DST-NRF Centre of Excellence in Strong Materials, hosted by University of the Witwatersrand, Johannesburg and School of Chemical and Metallurgical Engineering, University of the Witwatersrand, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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in the production mixture. For example, the product has to be novel, better, or cheaper than what is available elsewhere. This is where research is used to try and develop new capabilities, or to improve them, or to develop better production routes and methods. Thus the research is important, although it must be remembered that while ’blue skies’ research can lead to useful developments, the aim of the AMI is to remain focused on building local industry, so the research must have a perceived and needed impact. Figure 1 shows some of the applications for which PGMs have been used over the years, some being very long-standing, together with some of the areas that have been researched. Thus, this diagram represents a partial overview of some older applications, as well as some later developments. There is a wide variety of properties and products and it is the basis for a discussion of past projects and the identification of future research areas that could lead to new industries.

PGMs: A cornucopia of possible applications

and thermal conductivity of Pt are important (Fischer, 1992). Mechanically, Pt alloyed with rhodium or iridium combine high ductility with adequate creep strength and these properties give the alloys potential for applications in the chemical, space technology and glass industries (Fischer, 2001; Whalen, 1988; Lupton, 1990). In the spacecraft industry, PGMs are used to increase the heat resistance of rocket engine nozzles. A very important application is the manufacture of high-purity optical glasses and glass fibres by using platinum-containing tank furnaces, stirrers and feeders to withstand the high temperatures, mechanical loads and corrosive attack under the glass manufacturing and shaping conditions. In glasses, outstanding purity, homogeneity and the absence of bubble inclusions can be achieved only by using platinum alloys. If ceramic melting vessels were used, ceramic particles would be loosened by erosion, which would contaminate the glass melts and so compromise the desired optical properties such as transmittance. The alloys are expensive, but there are very few other materials that could replace platinum without risk of compromising the product. Thus, this is a valuable and safe use for platinum and some of its alloys.

Another ‘standard use’ of PGMs is in thermocouples and this arises because of their high electrical and thermal conductivity, as well as stability due to their high corrosion resistance. Typical alloys are Pt-Rh and Pt-Ru. The thermocouples do eventually degrade in aggressive environments and become brittle, but usually without much impairment of the electrical and conductive properties.

A long-standing use of PGMs is in catalysts and South Africa has established an autocatalyst manufacturing and refurbishment plant for automobiles. The alloys work very well and can be based on Pt or Pd, but are expensive. Competitors would like to replace these expensive alloys with something cheaper, but so far, nothing has been found that matches the good catalytic properties and robustness of the PGMs. There is also added flexibility because of work done

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with nanoparticles, which has the potential to increase the effectiveness of the catalysts, as the surface areas are increased. Much work has been done in South Africa on catalysts for different applications, both for PGMs, as well as for gold under the Autek project administered by Mintek. However, there is always a challenge with catalysis, because competitors are always looking for cheaper catalysts and there is a risk that the cheapness of other catalysts will offset the greater effectiveness of the PGM catalysts. The knowledge that palladium could be a ductile permeable storage medium for hydrogen dates from the 1970s (Knapton, 1977) and there is still work being done to optimise the properties, although it is used in hydrogen fuel cells.

Platinum has been used for jewellery for many years. Most of the alloys are 95 wt% Pt, with the other components usually being added to improve the hardness, since platinum is relatively soft. Various methods of surface hardening have been tried (Weber et al., 1996; McGill and Lucas, 1987). Other additions have been made to improve the flowability and hence castability. The most common alloying elements are copper, palladium, cobalt, gallium, iridium, ruthenium and indium. Copper is commonly added to make a generalpurpose alloy which casts well and is easy to work, and titanium has also been identified as a possible addition (Biggs et al., 2005). There is a 950 (95 wt%) hallmark for platinum alloys, which means that only 5 wt% is available for any additions. This hallmark is also synonymous with ’quality’ and closely guarded by the platinum industry. This situation is different from gold, where there are hallmarks for a wide range of alloys and gold contents. This means that there is much less flexibility in the platinum jewellery alloys, because they have to be 95 wt% Pt to be considered for hallmarking. Under the Innovation Fund, the Hot Platinum project was established. The aim was two-fold: to develop casting alloys for jewellery applications and also to develop small furnaces suitable for jewellers to use. Although some alloys were successfully developed, it was the furnaces that became more

PGMs: A cornucopia of possible applications

In 1911, Monnartz (1911) found that the corrosion of ironchromium alloys was much reduced by wrapping platinum wire about the samples, or by alloying with platinum. Subsequently, researchers added PGMs to various alloys to improve the corrosion resistance: Stern and Wissenberg (1959) added palladium to titanium; Tomashov (1958, 1967, 1970) added palladium to steels; Greene (1961) worked on chromium and high-chromium nickel-chromium alloys in sulphuric acids; Biefer (1970) added PGMs to type 430 stainless steel (and found Ru performed better than Pd); Tomashov (1976) added ruthenium to titanium and titanium-nickel alloys; Streicher (1977) added PGMs to stainless steels. Hoar (1960) identified platinum, palladium, rhodium and ruthenium as the best PGMs for cathodic medication, but as platinum and rhodium are expensive, ruthenium is likely to be the PGM of choice in South Africa. Green et al. (1961) gave a ranking of iridium, rhodium, ruthenium, platinum and palladium being more effective than osmium, gold or rhenium in maintaining a stable passivity. Although the ranking is not always obeyed, a discernable ranking of the alloying elements for improving corrosion is: Ir > Rh > Ru > Pt > Pd > Os > Au > Re, by comparing the corrosion rates and this was used by Potgieter (1990, 1991) in his discussion of the cathododic modification effect of the PGMs. Tomashov (1980) also looked at additions of 0.1 to 0.4 wt% Ru, Os, Ir, Pt or Pd and Chernova (1980) studied 0.2-0.4 wt% Ru. Tomashov et al. (1984) reported that Ru could block lattice defects, as well as enhancing the corrosion resistance in steels. Tomashov et al. (1969) added 0.1–0.4 wt% osmium, iridium and ruthenium to chromium alloys and found that the rate of dissolution in chromium-ruthenium alloys in the active state decreased with increasing Ru content. Tomashov et al. (1976) found that 0.2 %Ru addition improved the corrosion resistance of Ti-Ni alloys. Streicher (1977) found that Ru additions of 0.2 wt% conferred better corrosion resistance than the same amount of Pd. More recently, work in this area has increased, for example, Zhang. (2009) and Olaseinde et al. (2012), Potgieter (1995), to explain the effect, Van der Lingen and Sandenbergh (2001), and there is also work on ion implantation of PGMs in an attempt to retard stress corrosion cracking. Shing et al. (2001) added ruthenium to WC-Co alloys as a strengthener, with improved results and this also has potential for improving the corrosion resistance.

Nickel-based superalloys (NBSAs) have excellent mechanical properties due to precipitation strengthening of small, stable, ordered particles and have long been the ‘standard’ for aerospace alloys, although now lighter titanium alloys are being targeted for some of the components at lower temperatures. Platinum has similar chemistry to nickel and so reacts similarly with the alloying elements and has been investigated as a potential substitute in higher temperature alloys in even more aggressive environments (Coupland et al., 1980; Fischer et al., 1997, 1999a, 1999b, 2001; VĂślkl et

al., 2000; Wolff and Hill, 2000). The potential of using an fcc PGM analogue was mooted several times (Bard et al., 1994), in NIMS, Japan owing to the good properties of iridium-based and rhodium-based alloys (Yamabe et al., 1996, 1997, 1998a, 1998b, 1999; Yu et al., 2000), Pt-Ir alloys (YamabeMiterai et al., 2003), Pt-Al-Nb and Pt-Al-Ir-Nb alloys (Huang et al., 2004) and platinum-based alloys in South Africa (Wolff and Hill, 2000). The PGMs and nickel have similar structures (mostly fcc) and similar chemistries for the formation of similar phases. Advantages of PGM-based alloys over the NBSAs are the increased melting temperatures (e.g. 2443°C for iridium, 1769°C for platinum compared to 1455°C for nickel) and the excellent corrosion properties. Although platinum-based alloys would have been unlikely to replace all NBSAs on account of their higher price and higher density (Pt has a density of 21.5 g.cm-3, compared to nickel’s density of 8.9 g.cm-3), they were identified as having potential for use in components subjected to the highest temperatures. For Pt-based alloys, an increase in application temperature of at least 200°C could be gained (and more for Ir-based alloys). Although changes in engine design could be necessary, the higher application temperatures could offset the increased density and expense and the alloys could be recycled. Most PGMs have the face-centred cubic (fcc) structure, apart from ruthenium, which is hexagonal close-packed (hcp or cph). Iridium has a higher melting point than platinum, but has the disadvantage of brittleness (Panfilov et al., 2008) and is also in shorter supply. Thus, platinum is the preferred alloy base among the PGMs in the most extreme environments in terms of elevated temperatures, aggressive atmospheres and higher stresses (Wolff and Hill, 2000; Hill et al., 2001a; Cornish et al., 2003). Concerning coatings on these alloys, investigations have indicated that either no coatings would be necessary, or at least simpler coatings could be used than those currently used on nickel-based superalloys (Cornish et al., 2009a, 2009b; Douglas et al., 2009). Three major ranges of Pt-based alloys have been developed. One is based on Pt-Al-Cr-Ni (HĂźller et al., 2005; Wenderoth et al., 2005, 2007; Rudnik et al., 2008; VĂślkl et al., 2005, 2009) and two are based on Pt-Al-Cr-Ru (Cornish et al., 2009a, 2009b; Douglas et al., 2009). For the Pt-Al-CrRu alloys, one set was more malleable but less resistant to extreme chemical environments, whereas the other had higher chemical resistance, but was more difficult to form. The alloys showed very favourable mechanical properties (SĂźss et al., 2002). No alloys were produced commercially and the major problems was to identify a suitable application where the density of the alloys would not compromise their use, as they were too dense for the current designs of turbine engines. Another disadvantage is that they are extremely expensive. However, within these restrictions, the alloys could have potential as coatings on other lighter, more affordable substrates. Under the Platinum Development Initiative (PDI), much work was done to prove the potential of a Pt-based alloy and some alloys were developed that showed very good properties (Cornish et al., 2003, 2009a, 2009b; Cornish and Chown, 2011; Douglas et al., 2009; Potgieter et al., 2010), which were mainly targeted for land-based turbines and were based on Pt-Al-Cr-Ru in South Africa, with German colleagues in the collaboration preferring to base their alloys on Pt-Al-CrNi. The reason for their choice was a reluctance to use VOLUME 117

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successful and these are still being sold under the Hot Platinum brand, in large and small sizes. This is an example of a successful project.

PGMs: A cornucopia of possible applications ruthenium because of its ability to form RuO2, RuO4 and RuO3 (Caston, 1965) and it is the latter which has the highest partial vapour pressure, which was seen as a concern for long-term stability because of its volatility. The alloys performed well, but much better-supported parallel work outside of South Africa produced some lighter alloys which,.although they initially had stability problems, became better designed and were identified as having better potential in aerospace than the much denser Pt-based alloys, which would increase the momentum and hence the required strength in the moving parts of a turbine. Additionally, more work was done on the Ni-based superalloys which resulted in an increase the temperature range, although only by about 100°C. Most of the improved temperature range was due to the improved coatings (which provide a thermal barrier), as well as the design of turbines which employ forced cooling. The Pt-based alloys did have better properties at higher temperatures than the Ni-based superalloys and the volime fraction of the precipitates was increased during the course of the research to become closer to that of the NBSAs (Shongwe et al., 2010). The PDI work also allowed collaboration with Cambridge University, UK (Fairbank et al., 2000; Hill et al., 2001a) and NIMS, Japan (Hill et al., 2001b). However, this work was changed to focus on coatings and some interesting coatings with potential were developed. The bulk alloys remain and could be used for applications where density is not a problem, possibly in reactors for high temperature and corrosive environments. One of the interesting outomes from the work was the understanding of why there were two phase diagrams available for the base Pt-Al system (McAlister and Kahan, 1986; Oya et al., 1987). The diagrams differed in the shape of the platinum-based solid solution and well as the number and transformation temperatures of the Pt3Al phases. These differences were significant because they would affect the stabilities of the strengthening Pt3Al phases, as well as the amounts that could be produced. The German workers in the PDI collaboration (Fischer, 1992, 2001; Fischer et al. 1997 1999a, 1999b; VĂślkl et al., 2000; 2005, 2009; HĂźller et al., 2005; Rudnik et al., 2008; Vorberg et al., 2004, 2005; Wenderoth et al., 2007) preferred McAlister and Kahan’s (1986) diagram, while the South African workers preferred the McAlister and Kahan (1986) diagram for most of the phase boundaries, but Oya et al.’s (1987) diagram for the temperatures of transformation of the Pt3Al phases. The structure of one lower temperature phase was derived (Douglas et al., 2007) and showed some similarities to the modelled phase of Chauke et al. (2010). The only difference was the purity of the raw materials: those of the Germans were much purer (Cornish and Chown, 2011). This was an important observation, because the purity differences in the were small (less than 0.1 at.%), but the affect was significant. Later work (Odera et al., 2015) looked at reducing the Pt content without compromising the properties and substituted vanadium and/or niobium for some of the Pt, which could potentially increase the application temperature as well. An important outcome of this research was the development of a better method of etching the Pt-based alloys (Odera et al., 2012), which allowed the microstructure to be better resolved. This has improved the understanding of the microstructure-property relationship for different samples and

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has helped to determine the phase diagrams. Before the PDI work, Pt wires had been used to stabilize the position of components and Pt-Al coatings are used in the complex coatings of the nickel-based superalloys (Purvis and Warnes, 2001), because of their high temperature stability and strength. Some work continued on this using the experience derived from the bulk Pt-alloys.

Shape memory alloys can change their shape at a specific temperature and are used in many applications: sensors, temperature-sensitive switches, force actuators, fire safety valves, orthodontic wires, fasteners and couplers. The most well-known example is NiTi, commercially known as Nitanol. Biggs et al. (2003) showed that the addition of platinum increases the transition temperature of the alloy, which would allow different applications. By varying the Pt content, the transition temperature for the shape memory alloy can be varied between room temperature and 1000°C and so alloys could be developed for niche applications at specific temperatures. There is potential for adding a ternary component to the Pt3Al and PtTi phases and the PtFe3 phase has been recognized for its potential for low temperature applications. Work has been done at Mintek, the University of the Witwatersrand and the CSIR, but no alloy has been commercialized yet from South Africa. It is likely that it would be better to identify the niche, derive the required properties and then select the most likely alloy. PGM shape memory alloys could be used for dental applications, using the noble properties as well. In Germany, palladium is not permitted in dental alloys and so the target market must be assessed.

Magnetic materials are essential in many industries and there is a need to find substitutes for expensive rare earth metals. Historically, patents held by American and Japanese companies have been a barrier to entry to the magnet market, but China’s establishment of research institutes in the 1950s and 1960s and the expiry of the American and Japanese patents in the 1980s, has enabled China to enter and dominate the magnet market. Currently, China has a monopoly on the magnet value chain, producing 90% of the world’s permanent magnets. China also has most of the world’s rare earth resources. One of the new magnetic materials being investigated is the Fe-Co-Pd alloy, (Vokoun et al., 2005) which has potential applications in biotechnology, electronics and green technologies (hybrid vehicles and direct drive wind turbines). Magnetic thin films and nanostructures that are magnetized perpendicular to their surface are essential to many developing technologies, including spintronics devices (Mangin et al., 2006) and patterned media (Todorovic et al., 1999; Ding and Adeyeye, 2013), especially because of the need to maintain thermal stability as device dimensions are shrinking increasingly into the nanoscale. Multilayer or superlattice structures consisting of alternating ferromagnetic and non-magnetic layers are highly suitable for these applications due to their tunable perpendicular magnetic anisotropy (PMA) and saturation magnetization (Terris and Thomson, 2005; Rippard et al., 2010). By combining the soft magnetic properties of iron and the

PGMs: A cornucopia of possible applications

The PGMs have a wide range of uses. Some of these are well established, but others are new. Although the PGMs show excellent properties, because of their high cost there is always the risk of substitution, even by materials with slightly inferior properties. However, there are newer potential uses and these should be targeted, first by research and then if successful, more commercially. The additions of small amounts of PGMs to other materials is worth pursuing and probably one of the most exciting applications are magnets containing Pt and Pd.

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hard magnetic properties of cobalt and palladium, an optimal magnetic material could be developed (Vokoun et al., 2005). Fe-Co alloys have a B2 structure which makes them extremely brittle at room temperature (Matsuda et al., 2016). The workability of this alloy can be improved by substitution of nickel or palladium. Substitution of palladium leads to a remarkably high tensile strength and elongation at room temperature (Matsuda et al., 2016). Apart from the mechanical properties of the Fe-Co-Pd alloy and research on the magnetic shape memory properties, little work has been done on the ternary phase diagram. In order to understand the potential alloys better, the ternary phase diagram is needed, so that the compositions can be tailored to give the optimum phases for the desired properties.

PGMs: A cornucopia of possible applications ODERA, B.O., CORNISH, L.A., PAPO, M.J. and RADING, G.O. 2012. Electrolytic etching of platinum-aluminium based alloys. Platinum Metals Review, vol. 56, no. 4. pp. 257–261. ODERA, B.O., PAPO, M.J., COUPERTHWAITE, R., RADING, G.O., BILLING, D. and CORNISH, L.A. 2015. High-order additions to platinum-based alloys for high temperature applications. Journal of the Southern African Institute of Mining and Metallurgy, vol. 115. pp. 241–250. OLASEINDE, O.A., VAN DER MERWE, J.W., CORNISH, L.A. and OLUBAMBI, P.A. 2012. Electrochemical studies of Fe-21Cr-1Ni duplex stainless steels with 0.15wt% ruthenium at different temperatures. Journal of the Southern African Institute of Mining and Metallurgy, vol. 112, no. 7. pp. 535–538. OYA, Y., MISHIMA, U. and SUZUKI, T. 1987. L12 D0c martensitic transformation in Pt3Al and Pt3Ga. Zeitschrift für Metallkunde, vol. 78, no. 7. pp. 485–490.

TOMASHOV, N.D. ,KAZARIN, V.I., MIKHEEV, V.S. and GONCHARENKO, B.A. 1976. Effect of ruthenium on the electrochemical and corrosion behaviour of titanium and titanium-nickel alloys in acid chloride solutions. Protection of Metals, vol. 12, no. 5. pp. 471–474. TOMASHOV, N.D. 1980. Study of the potential oscillations of the alloys on the basis of chromium alloyed with the metals of the platinum group on solutions of sulphuric acid. Physical Chemistry Chemical Physics, vol. 84, no. 4. pp. 383–387. TOMASHOV, N.D., CHERNOVA, G.P. and USTINSKY, E.N. 1984. Effect of platinum elements additions on the active dissolution of plastic chromium in sulfuric acid. Corrosion, vol. 40, no. 3. pp. 134–137. VAN DER LINGEN, E. and SANDENBERGH, R.F. 2001. The cathodic modification behaviour of Ru additions to titanium in hydrochloric acid. Journal of Corrosion Science, vol. 43. pp. 577–590.

PANFILOV, P., PILUGIN, V.P. and ANTONOVA, O.V. 2008. On specific feature of plastic deformation in Ir. Creep 2008: Proceedings of the 11th International Conference on Creep and Fracture of Engineering Materials and Structures, Bayreuth, Germany, 4–9 May. Abstract CP-131.

VOKOUN, D., WANG, Y.W., GORYCZKA, T. and HU, C.T. 2005. Magnetostrictive and shape memory properties of Fe-Pd alloys with Co and Pt additions. Smart Materials and Structures, vol. 14, no. 5. S261 p.

POTGIETER, J.H., HEYNES, A.M. and SKINNER, W. 1990. Cathodic modification as a means improving the corrosion resistance of alloys. Journal of Applied Electrochemistry, vol. 20. pp. 711–715.

VÖLKL, R., FREUND, D., BEHRENDS, A., FISCHER, B., MERKER, J. and LUPTON, D. 2000. Platinum base alloys for high temperature space applications. Euromat 99 Series: Materials for Transport. Winkler, P.J. (ed.). Wiley-VCH, Weinheim.

POTGIETER, J.H. 1991 Alloys cathodically modified with noble metals. Journal of Applied Electrochemistry, vol. 21. pp. 471–482.

VÖLKL, R., YAMABE-MITARAI, Y., HUANG, C. and HARADA, H. 2005. Stabilizing the L12 structure of Pt3Al(r) in the Pt-Al-Sc system. Metallurgical and Materials Transactions A, vol. 36, no. 11. pp. 2881–2892.

POTGIETER, J.H., VAN BENNEKOM, A. and ELLIS, P. 1995. Investigation of the active dissolution behaviour of a 22% chromium duplex stainless steel with small ruthenium additions in sulphuric acid. ISIJ International, vol. 35. pp. 197–202. POTGIETER, J.H., MALEDI, N.B., SEPHTON, M. and CORNISH, L.A. 2010. The Platinum Development Initiative: platinum-based alloys for high temperature and special applications: Part IV - Corrosion. Platinum Metals Review, vol. 54, no. 2. pp. 112–119. PURVIS, A.L. and WARNES, B. M. 2001. The effects of platinum concentration on oxidation resistance of superalloys coated with single-phase platinum aluminide. Surface and Coatings Technology, vol. 146. pp. 1–6. RIPPARD, W.H., DEAC, A.M., PUFALL, M.R., SHAW, J.M., KELLER, M.W., RUSSEK, S.E., BAUERAND, G.E.W. and SERPICO, C. 2010. Spin-transfer dynamics in spin valves with out-of-plane magnetized CoNi free layers. Physical Reviews B. 81 p. RUDNIK, Y., VÖLKL, R., VORBERG, S. and GLATZEL, U. 2008. The effects of Ta additions on the phase compositions and high temperature properties of Pt base alloys. Materials Science and Engineering A, vol. 479. pp. 306–312. SHING, T.L., LUYCKX, S., NORTHROP, I.T. and WOLFF, I. 2001. The effect of ruthenium additions on the hardness, toughness and grain size of WC-Co. International Journal of Refractory Metals and Hard Materials, vol. 19. pp. 41–44. SHONGWE, M.B., ODERA, B., SAMAL, S., UKPONG, A.M., WATSON, A., SÜSS, R., CHOWN, L.H., RADING, G.O. and CORNISH, L.A. 2010. Assessment of microstructures in the development of Pt-based superalloys. Proceedings of Light Metals 2010, Muldersdrift, Johannesburg, 27–29 October 2010. Southern African Institute of Mining and Metallurgy, Johannesburg. pp. 184–202. STERN, M. and WISSENBERG, H. 1959. The influence of noble metal alloy additions on the electrochemical and corrosion behaviour of titanium. Journal of the Electrochemical Society, vol. 106, no. 9. pp. 759–764.

VÖLKL, R., WENDEROTH, M., PREUSSNER, J., VORBERG, S., FISCHER, B., YAMABEMITERAI, Y., HARADA, H. and GLATZEL, U. 2009. Development of a precipitation-strengthened Pt-base alloy. Materials Science and Engineering A, vol. 510-511. pp. 328-331. VORBERG, S., WENDEROTH, M., FISCHER, B., GLATZEL, U. and VÖLKL, R. 2004. Pt-AlCr-Ni superalloys: heat treatment and microstructure. Journal of Metals, vol. 56, no. 9. pp. 40–43. VORBERG, S., WENDEROTH, M., FISCHER, B., GLATZEL, U. and VÖLKL, R. 2005. A TEM investigationof the / ’ phase boundary in Pt-based superalloys. Journal of Metals, vol. 57, no. 3. pp. 49–51. WEBER, W., ZIMMERMANN, K. and BEYER, H.H. 1996. Surface-hardened objects of platinum and palladium and their method of production. US Patent 5,518,556. WENDEROTH, M., VÖLKL, R., VORBERG, S., YAMABE-MITERAI, Y., HARADA, H. and GLATZEL, U. 2007. Microstructure, oxidation resistance and high temperature strength of / ’ hardened Pt base alloys. Intermetallics, vol. 15. pp. 539–549. WENDEROTH, M., CORNISH, L.A., SÃœSS, R., VORBERG, S., FISCHER, B., GLATZEL, U. & Völki, R. (2005). On the Development and Investigation of Quaternary PtBased Superalloys with Ni Additions. Metallurgical and Materials Transactions, A, vol. 36 p.. WHALEN, M.V. 1988. Space station resistojets. Platinum Metals Review, vol. 32, no. 1. pp. 2–10. WOLFF, I.M. and HILL, P.J. 2000. Platinum metals-based intermetallics for hightemperature service. Platinum Metals Review, vol. 44, no. 4. pp. 158–166. YAMABE, Y., KOIZUMI, Y., MURAKAMI, H., RO, Y., MARUKO, T. and HARADA, H. 1996. Development of Ir-base refractory superalloys. Scripta Meterialia, vol. 35, no. 2. pp. 211–215.

STREICHER, M.A. 1977. Alloying stainless steels with the platinum metals. Platinum Metals Review, vol. 21. pp. 51–54.

YAMABE-MITARAI, Y., KOIZUMI, Y., MURAKAMI, H., RO, Y., MARUKO, T. and HARADA, H. 1997. Rh-base refractory superalloys for ultra-high temperature use. Scripta Meterialia, vol. 36, no. 4. pp. 393–398.

SÜSS, R., FREUND, D., VÖLKL, R., FISCHER, B., HILL, P.J., ELLIS, P. and WOLFF, I.M. 2002. The creep properties of Pt-base / ’ analogues to Ni-base superalloys. MaterialsScience and Engeering A, vol. 338. pp. 133–141.

YAMABE-MITARAI, Y., RO, Y., HARADA, H. and MARUKO, T. 1998a. Ir-base refractory superalloys for ultra-high temperature use. Metallurgical Transactions A, vol. 29, no. 2. pp. 537–549.

TERRIS, B.D. and THOMSON, T. 2005. Nano fabricated and self-assembled magnetic structures as data storage media. Journal of Physics, vol. D38. pp. 199–222.

YAMABE-MITARAI, Y., RO, Y., MARUKO, T. and HARADA, H. 1998b. Precipitation hardening of Ir-Nb and Ir-Zr alloys. Scripta Materialia, vol. 40, no. 1. pp. 109–115.

TODOROVIC, M., SCHULTZ, S., WONG, J. and SCHERER, A. 1999. Writing and reading of single magnetic domain per bit perpendicular patterned media. Applied Physics Letters, vol. 74. pp. 2516–2518.

YAMABE-MITARAI, Y., RO, Y., MARUKO, T. and HARADA, H. 1999. Microstructure dependence of strength of Ir-base refractory superalloys. Intermetallics, vol. 7, no. 1. pp. 49–58.

TOMASHOV, N.D. 1958. Methods of increasing the corrosion resistance of metal alloys. Corrosion, vol. 40. 229t.

YAMABE-MITARAI, Y. and AOKI, H. 2003. An assessment of Pt-Ir-Al alloys for high temperature materials. Journal of Alloys and Compounds, vol. 59. pp. 143–152.

TOMASHOV, N.D. 1967. Passivity and Protection of Metals Against Corrosion. 2nd edn. Springer. 83 p. TOMASHOV, N.D., CHERNOVA, G.P. and UTRINSKI, E.N. 1969. Platinum Metals Review, vol. 23, no. 4. pp. 143–149.

YU, X.H., YAMABE-MITARAI, Y., RO, Y. and HARADA, H. 2000. Design of quaternary Ir-Nb-Ni-Al refractory superalloys. Metallurgical and Materials Transactions A, vol. 31A, no. 1. pp. 173–178A.

TOMASHOV, N.D. CHERNOVA, G.P. and VOLKOV, L.N. 1970. Effect of palladium on corrosion and electrochemical behaviour of 0KH25N6T steel. Zaschita Metallov, vol. 6, no. 4. pp. 425–427.

ZHANG, L., ZHANG, W., JIANG, Y., DENG, B., SUN, D. and LI, J. 2009. Influence of annealing treatment on the corrosion resistance of lean duplex stainless steel 2101. Electrochimica Acta, vol. 54, no. 23. pp. 5387–5392. N

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Metal and alloy spheroidisation for the Advanced Metals Initiative of South Africa, using high-temperature radiofrequency plasmas by H. Bissett and I.J. van der Walt

Additive manufacturing presents an attractive and cost-effective method to produce complex designs, particularly when the material of construction includes precious metals such as Pt, Rh, Ir and their alloys. Layer deposition methods such as thermal spraying are also commonly used to apply protective coatings containing precious metals onto manufactured components. These methods require spherical powders to ensure a dense part or a defect-free layer. Thermal plasmas are suitable for spheroidisation of metal and alloy powders. A 15 kW plasma system from TEKNA Plasma Systems Inc. was purchased in 2016. The capability of this system was investigated by performing spheroidisation experiments making use of irregular-shaped titanium metal powder at various plasma operating conditions. The resulting powders were characterised in terms of morphology, density and flowability. The flow characteristics of the powder were determined by means of a Hall flow test. The plasma treatment resulted in an increase in spheroidisation ratio and fraction of evaporation with increasing plasma plate power. The treated powder displayed improved flow characteristics.

7 &23,. additive manufacturing, laser sintering, powder metallurgy, spheroidisation, titanium, radio-frequency plasma.

0532,'-5420 Titanium and its alloys show good corrosion resistance due to the stability of the oxide films that form spontaneously on the surface. The addition of precious metals such as Pd, Ru and Pt can further increase the corrosion resistance of metals and alloys, including alloys systems such as Fe-Al and stainless steels (Mwamba, Cornish and van der Lingen, 2014; Couperthwaite, 2015; Olubambi, Potgieter and Cornish, 2009), but at increasing cost. In recent years the manufacturing landscape has changed significantly with an emphasis on developing technologies to manufacture components cost- and resourceefficiently on a small scale (Williams and Revington, 2010), making this technology suitable for expensive powders containing precious metals. The emergence of additive manufacturing (AM) heralds a technology where value chains are shorter, smaller, localised and more collaborative (Ford and Despeisse, 2016). One of these new technologies used to manufacture high-quality components is direct laser sintering, or 3D

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* The South African Nuclear Energy Corporation (SOC) Ltd. (Necsa), Pretoria, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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printing. The powders required for these methods have to be spherical and of a narrow particle size distribution as this affects the packing density and sintering mechanism. The powder also has to be chemically pure, as impurities such as H, O, C, N and S causes brittleness, influence metal properties such as tensile strength, hardness and ductility and also increase surface tension during processing (Despa and Gheorghe, 2011). More conventional techniques which can also reduce cost where materials are expensive include layer deposition methods such as thermal spraying (Wrona et al., 2017). Thermal spraying methods such as plasma spraying, however, require a powder with excellent flowability and high apparent density to produce dense coatings (Qian, Du and Zhang, 2009). In this regard, spherical powders present optimum flowability and density. Thermal plasmas, characterised by their extremely high temperatures (3000–10 000 K) and rapid heating and cooling rates (approx. 106 K/s) under oxidizing, reducing or inert conditions, are particularly suitable for spheroidisation of metal and alloy powders with relatively high melting points. Residence times of particles in the plasma region range from 5 to 20 ms, but this is usually sufficient as 7 to 8 ms is required for heating and melting of metal particles in the 30–50 Οm size range at 3500 K. A radio-frequency (RF) plasma is the preferred plasma method due to the longer residence time, larger plasma volume and the lower risk of contamination as no electrode erosion occurs (Grignard, 1998). TEKNA Plasma Systems Inc. is one of the companies that supplies high-temperature plasma technology equipment for various applications, including plasma spheroidisation.

Metal and alloy spheroidisation for the Advanced Metals Initiative of South Africa Various research studies utilizing TEKNA equipment have been conducted. Spheroidisation of ceramics such as Al2O3, Cr2O3, SiO2, ZrO2, WC and TiC, as well as W and Mo metal powders have been studied making use of various TEKNAtype torches and various sized systems (Gignard, 1998; Jiang and Boulos, 2006; Li and Ishigaki, 2001). From available literature it is clear that TEKNA is one of the leaders in plasma spheroidisation technology for high-melting-point materials. The South African Nuclear Energy Corporation SOC (Necsa) has recently purchased a 15 kW RF thermal plasma system from TEKNA Plasma Systems through the National Equipment Programme, managed by the National Research Foundation of South Africa. In early 2017 the system, especially designed for metal powder spheroidisation, was commissioned. At the time of writing this article, work relating to the Precious Metals Development Network (PMDN), Light Metal Development Network (LMDN) and the Nuclear Materials Development Network (NMDN) had already commenced and yielded promising results with materials such as Ti, Fe and ZrC. The focus of this study is the spheroidisation of irregularly shaped pure titanium metal powder to investigate the capabilities of the 15 kW plasma system and to introduce the system to prospective users requiring spheroidized ceramic, metal or alloy powders. The processing characteristics for the spherical Ti particle formation were investigated at various processing conditions. The particles obtained were characterised in terms of morphology, density and flowability. Spheroidisation of powders should result in an optimum spheroidisation ratio (%), while minimizing the fraction of material evaporated. Although the current study focused on titanium metal, the data obtained can be used to evaluate the spheroidisation potential for PGM and other alloy powders that are applicable to all the networks of the Advanced Metal Initiative (AMI).

influence the time a particle will remain in the plasma flame (residence time). The characteristics of the induction plasma and the degree of spheroidisation are dependent on the operating conditions such as the carrier gas flow rate, powder feed rate, central and sheath gas flow rate, sheath gas composition, plasma plate power and reactor pressure.

% % % % $ $% !" TEKNA’s 15 kW induction plasma system (Figure 2) makes use of a PL-35M induction torch which can operate between 2 and 5 MHz. A simple schematic representation of the system is shown in Figure 3. The plasma system is enclosed in a Faraday cage to prevent release of high-frequency (hf) radiation to the environment. Due to the fact that H2 is used in some instances, the system safety protocols require that the system is leak-tight, eliminating any danger of explosion. The system is intended to be used to process batch samples of powders up to 3 kg, depending on the initial bulk density

1734+7056/ $ # %# !# % $ $%! % " # For an induction torch as shown in Figure 1, the energy coupling between the high-frequency generator and the plasma is achieved by a cylindrical induction coil. The absence of electrodes eliminates the possibility of product contamination and allows for the use of various gases during plasma treatment. The plasma torch essentially consists of two concentric tubes with a small annular gap between them. The outer tube is known as the plasma-confinement tube and is water-cooled. The intermediate tube, usually made from quartz, extends down to the first turn of the induction coil. It serves to induce a flow pattern in the torch with a relative high-velocity sheath gas. The sheath gas in most instances is a mixture of Ar and H2 or Ar and He. The addition of H2 or He reduces heat transfer to the water-cooled outer tube due to the fact that the ionisation potentials of these gases are lower than that of Ar (Boulos, 1985). The design of the plasma torch allows for the investigation of various parameters in order to optimise the powder spheroidisation process. The powder to be treated is fed through a height-adjustable, water cooled probe into the plasma flame making use of a carrier gas (Ar) which will

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of the powder. This unit is not designed for, nor suitable for, production of highly pyrophoric materials such as Al, Mg, B and Zr. The system is equipped with several safety devices such as a rupture disk to prevent any potential overpressures and to allow for safe operation. The powder is introduced into the feeding probe via a vibration feeder and swept into the plasma by the carrier gas. For the purpose of this study the plasma torch was mounted on a reactor chamber equipped with a catch-pot for the collection of the solid particles. During plasma treatment evaporation of the particles occurred, resulting in the formation of fine deposits that were collected at the bottom of the cyclone and the filter. After completion of plasma treatment, the system was allowed to cool before the powder was passivated, to ensure safe removal.

" # " !$ % "! For all experiments at least 5 g of Ti powder was fed into the plasma at a flow rate of approximately 0.22 kg/h using a carrier gas (Ar) flow rate of 2 standard litres per minute (slpm) with the sheath gas containing 7.5 v/v% H2. Experiments were conducted at plasma plate powers of 9, 11 and 13 kW. The degree of evaporation was evaluated from the weight of fine deposits expressed as a ratio of the total weight of powder collected. In order to evaluate flowability of the powder before and after plasma treatment, in a single run 80 g of titanium powder was treated at 13 kW to obtain at least 50 g of densified powder. Pure titanium metal powder in the size range 150–180 Οm was used as feed material for all experiments. To ensure that this specific size range was present, the powder was sieved using 150 and 180 Οm (100 – 80 mesh) sieves.

optical microscope equipped with a Nikon D70s camera. The ‘spheroidisation ratio’, denoted as the ratio of the number of spherical particles to the total number of particles in the treated product, was estimated by counting the spherical particles within a specific frame size of the optical micrograph. Partly melted or elliptical-shape particles were not counted as spherical particles. The feed and the plasma-treated powders were also characterised according to density, using an AccuPyc II 1340 gas displacement helium pycnometer. The 1 cm3 sample cup was used for density analysis of three samples to calculate the average density and standard deviation. To investigate the internal structure of the treated Ti particles, powder products were embedded in Pratley Steel QuicksetŽ instead of the epoxy cold-mounting resin typically used in metallographic sample preparation method. The sample was polished to expose the cross-sections of the particles for examination by scanning electron microscopy (SEM, Quanta 200 3D with an Apollo X silicon drift detector). A SEM image of the surface of plasma treated particles was also obtained. As indicated earlier, a large amount of powder was plasma-treated at 13 kW to obtain at least 50 g of densified powder to be characterised according to flowability. The feed powder and the plasma-treated powder were characterised using an ASTM International Hall flow test method for flow rate of powders. This method necessitates that 50.0 g of metal powder is loaded into a funnel with a specified size orifice at the bottom. Once the powder starts flowing, T0 is noted and once all the powder has flowed through the orifice the end time is noted. If a powder does not flow through a 2.5 mm orifice, then a 5 mm orifice is used. Obviously, the shorter the time recorded, the better the flow properties of the powder (ASTM International).

97.'/5.860,8,4.-'..420 "" % " Figure 4 shows an optical micrograph of the irregularly shaped pure titanium powder. This powder was sieved to ensure that the size range 150–180 Οm was used as feed material. The density of the feed powder was 4.3107 ¹ 0.0023 g/cm3.

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The feed powder and the powder collected in the catch-pot after each run were characterised by using a Nikon SMZ-T2

Metal and alloy spheroidisation for the Advanced Metals Initiative of South Africa " # # $!# %$ % $ !#$ %" $ $!# The experimental conditions for spheroidisation are shown in Table I. The plasma-treated Ti powders were collected in the catch-pot of the reactor chamber, while the fine deposits were collected from the reactor chamber as well as from the collections pots attached to the cyclone and the filter. The total weight of powder was calculated as the sum of the densified powder and the fine deposits. This weight was then used to calculate the fraction of evaporation (%). Figure 5 illustrates the results of the experiments conducted at 9, 11 and 13 kW plasma plate powers. The treated powders obtained at higher plasma plate powers had a higher ratio of spherical particles in the processed products. Almost complete spheroidisation was obtained at 13 kW. It was, however, also observed that an increase in spheroidisation ratio resulted in an increased fraction of evaporation, ranging from 5.50 to 25.05 wt%. A higher fraction of evaporation equates to a loss of potential product or spherical particles due to the fact that more fine particles are collected in the cyclone and filter sections, as indicated in Figure 3. A high fraction of evaporation can be detrimental to the spheroidisation process for expensive materials such as Pt and Pd, or alloys containing these elements. Although the material is not lost, but condensed as fine particles, this can still make the process uneconomical,

especially if no use can be found for the fine particles. In this regard it might be preferable to perform the plasma treatment at a lower plasma plate power, yielding a lower spheroidisation ratio and lower fraction of evaporation and then repeating the exercise in order to increase the spheroidisation ratio while maintaining a low fraction of evaporation. Figure 6 illustrates the density and Figure 7 the morphologies, of the plasma-treated powder as a function of the plasma plate power. It can be seen that a slight densification of the feed powder occurred at 9 kW plasma plate power, with the density increasing from 4.3107 to 4.3177 g/cm3. A large amount of irregular shaped particles were still present, as shown in the optical micrograph Table I

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Metal and alloy spheroidisation for the Advanced Metals Initiative of South Africa (Figure 7). Increasing the plasma plate power to 11 kW had a dramatic influence on the density, which increased to 4.4078 Âą 0.0024 g/cm3. At 13 kW a slight decrease in the density, to 4.3870 Âą 0.0024 g/cm3, of the plasma-treated powder was observed, although more spherical particles were obtained according to the optical micrograph. The slight decrease in density might indicate the creation of internal voids within the particles, which can cause the formation of structural defects during manufacturing of components by AM methods such as 3D printing. A high degree of spheroidicity and smooth surface were apparent for all the spherical particles, as shown in Figure 7 by the artifact seen on top of all the spherical particle surfaces caused by the reflection of the microscope light. A key parameter for spheroidisation is heat transfer from the plasma discharge to the particles. This is determined by the plasma temperature and the residence time of the particles. The residence times for particles can be different, as not all particles will travel the same pathways and the particles will therefore be exposed to different temperatures within the plasma. The dynamics of mass transfer determine the time for complete melting. The thermophysical properties involved are molten liquid density, surface tension and viscosity. From the results above it can be seen that the particles treated at higher plasma plate powers were exposed to higher temperatures if it is assumed that the residence times of the particles were similar, resulting in high spheroidisation ratios at higher applied plate powers. Considering the slight decrease in density as the plasma plate power increased from 11 to 13 kW, the following explanation can be given. At a high plasma plate power (temperature) it can be assumed that the particles will be completely melted. At high quench rates (as is the case with plasma processing), the outer shells of the particles will solidify while the centres

remains in the liquid state and at a lower density, forming a cavity during complete cooling of the particles. Treating large particles might result in the formation of enlarged cavities (Grignard, 1998).

$ !# "% # ! ! " The surface and the internal structure of the plasma-treated Ti particles were investigated by SEM. The surfaces of the spherical particles were smooth, as shown in Figure 8 for the spherical particles obtained at a plasma plate power of 13 kW. The surfaces of the spherical particles observed at 9 and 11 kW were similar. The size of the particles was around 160 Îźm. The exposed cross-sections of three spherical particles are shown in Figure 9. The particles were either fully dense internally (left), contained voids (middle) or completely hollow (right). These flaws were obviously introduced during the melting and solidification processes discussed earlier regarding the formation of a solid outer shell and a liquid centre. Therefore it can be assumed that the particle in Figure 9 (right) was melted completely and rapidly quenched, while the particle in Figure 9 (left) was melted completely and then quenched at a slower rate, allowing the centre and shell to solidify at similar rates, resulting in optimum densification. From the results in Figure 6 it can be deduced that more of the spherical particles obtained at 13 kW contained cavities or were hollow and not fully densified, considering that the theoretical density of titanium is 4.506 g/cm3. Although the spheroidisation ratio at 11 kW was lower than at 13 kW, it can be assumed that more fully densified particles will be present in this sample. In order to determine what fraction of particles will be completely dense at a given condition, more experimental data and powder analysis will be required.

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Metal and alloy spheroidisation for the Advanced Metals Initiative of South Africa

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" % $ # #! The Hall flow tests were conducted for flowability determination on 50 g of feed powder as well as 50 g of densified product powder after plasma treatment at 13 kW. The feed powder was initially tested using a Hall flow funnel with an orifice size of 5 mm, but the time required for all the powder to pass through was less than 37 seconds and therefore only the 2.5 mm orifice was used to test the powder flowability. Three measurements were taken to obtain an average and standard deviation. The time required for 50 g of feed powder to flow through the 2.5 mm orifice was 102.9 Âą 1.0 seconds and 34.8 Âą 0.4 seconds for the densified powder. This result shows a significant improvement in the flowability of the powder after plasma treatment and clearly indicates that the plasma treatment increased the density and the spheroidicity of the particles.

20-/'.420. Highly spherical titanium particles were produced by thermal plasma treatment of irregular shaped titanium particles under argon-hydrogen thermal plasma in the TEKNA plasma system. The spheroidisation ratio and the fraction of evaporation increased from 35 to 90% and 5.5 to 25.05% respectively for plasma plate powers at 9 and 13 kW. A maximum powder density of 4.41 g/cm3 was obtained at 11 kW. The slight decrease in powder density at 13 kW was due to the presence of some hollow particles in this sample. This was attributed to the rapid quenching of particles causing a difference in the density of the solidified particle outer shell and the molten centre. During cooling a hollow particle was formed. The Hall flow test showed that the flow characteristic of the plasma-treated titanium powder was improved significantly, indicating that spheroidisation and densification were successfully carried out using the TEKNA plasma system. The results in this study show that spheroidisation of titanium metal is possible and therefore this system can be employed for the spheroidisation of other metal and alloy powder, including alloys containing precious metals.

- 02&/7,(7+705. The author would like to acknowledge the National Research Foundation and the South African Nuclear Energy Corporation for the financial support in purchasing the spheroidisation equipment. Mr Ryno van der Merwe, research scientist at Necsa, is thanked for the SEM analysis of the

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spherical particles. The Department of Science and Technology (DST) and the Advanced Metals Initiative (AMI) are also acknowledged for their financial support and for the opportunity to present this work. Necsa wishes to thank TEKNA Plasma Systems Inc. for the successful installation and commissioning of this equipment and the training they gave in this regard.

97*7370-7. ASTM INTERNATIONAL. (2006). Standard test method for apparent density of free-flowing metal powders using the Hall flowmeter funnel. Designation: B212 99. West Conshohocken, PA. BOULOS, M.I. (1985). The inductively coupled R.F. (radio frequency) plasma. Pure & Applied Chemistry, vol. 57, no. 9. pp. 1321–1352. COUPERTHWAITE, R.A. (2015). Effect of processing route on the microstructure and properties of an Fe-Al alloy with additions of precious metal. Materials Today. Proceedings 2, pp. 3932–3942. DESPA, V. and GHEORGHE, I.G. (2011). Study of selective laser sintering: A qualitative and objective approach. Scientific Bulletin of Valahia University Materials and Mechanics, vol. 6. pp. 150–155. DIGNARD, N.M. (1998). Experimental optimization of the spheroidization of metallic and ceramic powders with induction plasma. PhD thesis, Universite de Sherbrooke, Sherbrooke, Quebec, Canada. FORD, S. and DESPEISSE, M. (2016). Additive manufacturing and sustainability: an exploratory study of the advantages and challenges. Journal of Cleaner Production, vol. 137. pp. 1573–1587. JIANG, X. and BOULOS, M. (2006). Induction plasma spheroidization of tungsten and molybdenum powder. Transactions of the Nonferrous Metal Society of China, vol. 16. pp. 13–17. LI, Y. and ISHIGAKI, T. (2001). Spheroidization of titanium carbide powders by induction thermal plasma processing. Journal of the American Ceramic Society, vol. 84, no. 9. pp. 1929–1936. MWAMBA, I.A., CORNISH, L.A. and VAN DER LINGEN, E. (2014). Effect of platinum group metal addition on microstructure and corrosion behaviour of Ti47.5 at-%Al. University of Pretoria Institutional Repository, Pretoria, South Africa. http://hdl.handle.net/2263/41034 OLUBAMBI, P.A., POTGIETER, J.H. and CORNISH, L. (2009). Corrosion behavior of superferritic stainless steels cathodically modified with minor additions of ruthenium in sulphuric and hydrochloric acids. Materials and Design, vol. 30. pp. 1451–1457. QAIN, Y., DU, L. and ZHANG, W. (2009). Preparation of spherical Y2SiO5 powders for thermal-spray coating. Particuology, vol. 7. pp. 368–372. WILLIAMS, J.V. and REVINGTON, P.J. (2010). Novel use of an aerospace selective laser sintering machine for rapid prototyping of an orbital blowout fracture. International Journal of Oral and Maxillofacial Surgeons, vol. 39. pp.182–184. DOI: 10.1016/j.ijom.2009.12.002 WRONA, A., BILEWSKA, K., LIS, M., KAMIĹƒSKA, M., OLSZEWSKI, T., PAJZDERSKI, P., WIECĹ AW, G., JAĹšKIEWICZ, M. and KAMYSZ, W. (2017). Antimicrobial properties of protective coatings produced by plasma. Surface and Coatings Technology, vol. 318. pp. 332–340. N

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a9

A comparative study of hardness of PtCr-V and Pt-Al-V alloys in the as-cast and annealed conditions by B.O. Odera*†‥ and L.A. Cornish†‥#

Microhardness values of alloys in the Pt-Cr-V and Pt-Al-V systems were determined using an indenter which incorporated an optical microscope. The values were analysed with respect to the phases identified in each of the alloys. Identification of the phases had been done earlier using scanning electron microscopy with energy dispersive X-ray spectroscopy and X-ray diffraction. The hardness values were determined for samples in the as-cast and annealed conditions and superposed on the solidification projection and isothermal section at 1000°C. The results showed that the hardness depended largely on the phases present in the alloys and generally increased after annealing. However, in a few cases, grain growth and the resulting microstructural coarseness resulted in lower hardness values after annealing. The hardness of the alloys of the Pt-Cr-V and PtAl-V systems was also compared with those of Pt-Al-Cr and other Pt-Al based alloys. Comparison was also made with some Pt-Al based alloys of the higher order systems such as Pt-Al-Cr-Ru-V and Pt-Al-Cr-Ru-V-Nb. In general, higher hardness was exhibited by alloys containing ternary phases. 72":0-+3 Pt-based ternary alloys, Pr-Cr-V, Pt-Al-V, hardness, microstructure.

1.-0+ ).,01 Platinum-based alloys for high-temperature corrosive environments have been investigated as possible replacements for nickel-based superalloys (NBSAs) in the hotter parts of gas turbines (Wolff and Hill, 2000; Hill et al., 2001b). The NBSAs have excellent mechanical properties due to their γ/γ' structure but are limited by their maximum operating temperatures. Currently, the maximum temperature at which NBSAs operate is approximately 1100°C, which is approximately 85% of their melting temperature (Sims, Stoloff and Hagel, 1987). There is a need to develop new alloys that can operate at higher temperatures. The advantages include increased efficiency, which is related to the combustion temperature and higher temperatures would enable greater thrust, improved fuel efficiency and reduced pollution. Platinum has the same fcc structure and similar chemistry to nickel, as well as a higher melting point (1769°C for platinum, compared to 1455°C for nickel) and better corrosion resistance than nickel (Sßss et al., 2003).

VOLUME 117

* Department of Mechanical Engineering, Cape Peninsula University of Technology, Cape Town. †School of Chemical and Metallurgical Engineering, University of the Witwatersrand, Johannesburg. ‥ African Materials Science and Engineering Network (AMSEN), a Carnegie-IAS Network. # DST-NRF Centre of Excellence in Strong Materials, hosted by University of the Witwatersrand, Johannesburg. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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Alloy development work done by other researchers demonstrated that ternary and quaternary Pt-based alloys can exhibit a γ/γ′ two-phase structure, similar to nickel-based superalloys (Hill et al., 2001; Cornish Fischer and VÜlkl, 2003). Although the use of Ptbased alloys as a replacement for Ni-based superalloys is limited due to their higher price and higher density, it is possible that they could be used for critical components, or as corrosion-resistant coatings (Sßss et al., 2009). Some of the quaternary alloys exhibiting the γ/γ′ two-phase structure were of the PtAl-Cr-Ru system. Alloys of the quinary Pt-AlCr-Ru-V system have also been investigated as potential replacement of NBSAs (Odera, 2013; Odera et al., 2015). Vanadium replaced some of the Pt to reduce density and cost, as well as acting as a solid solution strengthener. Pt-Al-V and Pt-Cr-V are two component ternary systems of the quinary and a study of the microstructures of selected alloys of the systems and their ternary phase diagrams formed part of the investigation (Odera et al., 2012b; Odera, 2013; Odera et al., 2014). This paper presents hardness values of selected alloys of the two ternary systems in the ascast and annealed conditions and compares and discusses these values in relation to the microstructures. There is an empirical relationship between hardness and strength of a metal. Since hardness testing is easier to perform than

A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys tensile strength testing, especially for small samples, it is often used to obtain an indication of the strength of a metal, especially when only small samples are available (Dieter, 1986; Smith, 1990).

69#2-, 21.4/5#-0)2+ -2 Twenty-four alloy buttons weighing approximately 2 g each were prepared from Pt, Al, Cr of 99.9% purity and V of 99.6% purity. Out of these, 11 samples identified as Alloy 1 to Alloy 11 were of the Pt-Al-V system, while those identified as Alloy 12 to Alloy 24 were of the Pt-Cr-V system, all in the as-cast condition. The annealed samples were identified as 1H to 11H and 12H to 24H respectively. The samples were manufactured by arc-melting under argon atmosphere on a water-cooled copper hearth with a Ti oxygen-getter. Each button was turned over and re-melted three times in an attempt to achieve homogeneity. The samples were then halved and one half of each was metallographically prepared in the as-cast condition, while the other half was annealed for 1500 hours before preparation. The samples were ground on silicon carbide down to 1200 grit and then diamond polished down to 1 Οm. The samples were then etched in a solution of 10 g NaCl in 100 cm3 HCl (32% vol. concentration) (Odera et al., 2012a). Etching was done in a fume cupboard using a DC power supply and a voltage range of 9 V to 12 V gave adequate results. The current density in the electrolyte was approximately 100 A.m–2. The counter-electrode was a stainless steel wire suspended in the electrolyte solution. The indenter used incorporated an optical microscope with a maximum magnification of 1000 times and only an etched surface could be seen clearly in the machine. In the unetched condition, it was difficult to differentiate between the sample and the mounting resin surface. The hardness of the alloys was measured using a Vickers microhardness tester with a load of 300 g. At least

five measurements were carried out to obtain an average hardness value.

;23 /.3 Figure 1a shows an indentation on as-cast Alloy 5, average composition Pt83.9:Al6:V10.1 (at.%) and Figure 1b shows an indentation on annealed Alloy 5H of the same average composition. The indentations on as-cast Alloy 24 and annealed Alloy 24H, average composition Pt34.1:Cr57.4:V8.5 are shown in Figure 2. Figures 3 and 4 show indentations on as-cast Alloys 15 and 18, average compositions Pt53.0:Cr22.4:V24.5 and Pt25.8:Cr42.9:V31.3 respectivley. Table I gives hardness values (HV0.3) for as-cast and annealed PtAl-V alloys, together with the phases in each of the alloys, while Table II gives the same data for Pt-Cr-V alloys.

,3) 33,01 Alloys with a Pt content more than 70 at.% had the lowest hardness values (less than 500 HV0.3), while 64% of the alloys had relatively high hardness values of more than 500 HV0.3. The single-phase ternary Alloy 11 had the highest hardness of 886 HV0.3, which would be expected and is likely to be brittle. Generally, the load of 300 g used was too low to cause substantial cracking during indentation on these particular alloys. However, Alloy 11 exhibited cracks before indentation, confirming its brittle nature. Hill et al. (2001a, 2001b) reported Vickers macrohardness measurements on several Pt-based ternary alloys at room temperature. Eight of these were Pt-Al-Z alloys (where Z was Ni, Ru or Re). Six of the alloys had Vickers hardness less than 500 HV in the as-cast and annealed conditions and the highest value was 530 HV (the load used was not specified). In his investigation of Pt-Al-Cr alloys, SĂźss (2007) found that 75% of the alloys in the as-

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A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys

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Table I

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'43235#-2321.5 ,15.'2541124/2+54//0"3

1

Pt64.3:Al26.6:V9.1

655±27

635±13

~Pt3Al, ~Pt5Al3, ~Pt2Al, ~Pt3V, PtAl

~Pt3Al, Pt5Al3, ~Pt2V, PtAl

2

Pt59.1:Al23.1:V17.8

746±35

624±26

~Pt2V, ~Pt5Al3, PtAl

~PtV, ~Pt5Al3

3

Pt69.3:Al9.6:V21.1

519±34

516±17

~Pt3V, ~Pt3Al, ~Pt2Al

~Pt2V, ~Pt3Al, ~PtV

4

Pt69.8:Al22.3:V7.9

671±23

696±20

~Pt3Al, ~Pt2Al

~Pt3Al, ~Pt2Al

5

Pt83.9:Al6:V10.1

420±20

482±14

(Pt), ~Pt3Al

(Pt), ~Pt2V, ~Pt3Al

6

Pt80.9:Al4.3:V14.8

381±13

378±16

(Pt), ~Pt3Al

(Pt), ~Pt3Al

7

Pt52.5:Al22.6:V24.9

767±37

733±18

~PtV, ~PtV3, ~Pt5Al3

~PtV, ~PtV3, ~Pt5Al3

8

Pt53.7:Al8.8:V37.5

584±19

585±19

~PtV, ~PtV3, ~Pt5Al3

~PtV, ~PtV3, ~Pt5Al3

9

Pt84:Al2:V4

475±9

347±10

(Pt), ~Pt3Al

(Pt), ~Pt3Al

10

Pt72.4:Al13.7:V13.9

462±14

535±23

~Pt3Al, ~Pt2Al

~Pt3Al

11

Pt56.9:Al25.4:V17.7

886±23

-

Ï„2

~PtAl, ~PtV

Table II

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12

Pt73.8:Cr16.9:V9.3

272±8

280±26

~Pt3V

~Pt3V

13

Pt72.3:Cr8.3:V19.4

317±8

330±17

~Pt3V

~Pt2V, Pt3V

14 Ï„3

Pt33.4:Cr45.8:V20.8

527±28

589±23

~Cr3Pt, A15(~Cr3Pt), Ï„3

15

Pt53.0:Cr22.4:V24.5

936±59

637±28

~CrPt, Ï„3

16

Pt68.5:Cr16.2:V15.3

291±19

350±8

~Pt3V

~Pt3V

17

Pt32.8:Cr28.6:V38.6

854±30

1086±46

~Cr3Pt, Ï„3

~Pt2V, Ï„, ~CrPt

18

Pt25.8:Cr42.9:V31.3

1109±25

-

Ï„3, A15(~Cr3Pt), ~Cr3Pt

A15(~Cr3Pt), ~Cr3Pt, Ï„3

19

Pt38.8:Cr14.2:V47.0

796±35

923±57

~PtV3, ~PtV, Ï„3

~PtV3, ~PtV, Ï„3

20

Pt16.8:Cr13.4:V69.8

618±25

-

(V,Cr), ~PtV3

(V,Cr), ~PtV3

21

Pt28.7:Cr27.2:V44.1

1086±28

-

~Cr3Pt, ~PtV3, Ï„3, (V,Cr)

~PtV3 Ï„, (V,Cr)

22

Pt16.5:Cr56.1:V27.4

707±41

837±46

(V,Cr), A15(~Cr3Pt)

(V,Cr), A15(~Cr3Pt)

23

Pt14.9:Cr73.7:V11.4

692±73

544±20

(V,Cr), A15(~Cr3Pt)

(V,Cr), A15(~Cr3Pt)

24

Pt34.1:Cr57.4:V8.5

358±9

838±27

~Cr3Pt, A15(~Cr3Pt)

~Cr3Pt, A15(~Cr3Pt)

cast condition had Vickers hardness over 600 HV10 at room temperature. Figure 5 shows hardness values of as-cast Pt-Al-V alloys superimposed on the partial solidification projection at the Pt

'43235#-2321.5 ,15.'2541124/2+54//0"3

~Cr3Pt, ~CrPt A15(~Cr3Pt), ~Cr3Pt, Ï„3

rich corner derived and reported by Odera (2013). It shows the hardness values increasing away from the Pt-rich corner. The composition of the two ternary phases Ï„1 and Ï„2 are approximately V24Pt56Al20 and V18Pt57Al25 respectively. VOLUME 117

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A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys

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The hardness values for annealed alloys of the Pt-Al-V system are given in Table III. Seventy per cent of the annealed alloys had hardness values higher than 500 HV and the rest, which had relatively high Pt contents above 80 at.%, had hardness values below 500 HV0.3. About 45% of the alloys had hardness values over 600 HV0.3. Annealed Alloy 11H cracked and disintegrated while being ground in preparation for polishing and etching, indicating that it was very brittle and it also contained the ternary phase in as thecast condition. In all of the alloys, there was no cracking or noticeable deformation around the indentations, therefore their toughness could not be quantitatively assessed. However, the absence of notable cracking suggests reasonable toughness (except for Alloy 11 which comprised the ternary phase in the as-cast condition and had cracks before indentation). However, this could also be attributed to the relatively small load used in the microhardness testing. The small load was chosen because of the small size of the samples (approx. 1 g). However, these hardness values for the alloys, whether in the as-cast or annealed condition, are very high compared to that of pure platinum, which is very soft and ductile with a Vickers hardness of 50 HV (the load was not specified) (Murakami, 2008). These results are comparable to the findings of SĂźss (2007) in his investigation of Pt-Al-Cr alloys, where just over half of the annealed alloys had hardness values above 600 HV10. Only one of the Pt-Al based alloys, Pt62:Al19:Ni19, investigated by Hill (2001) had a hardness value of 530 HV (the load used was not specified). The rest had hardnesses below 500 HV. The phase ~Pt2Al disappeared during annealing of Alloy 1H, but the hardness remained more or less the same. In Alloy 2H, two phases, ~Pt2V and ~PtAl, disappeared and a new phase, ~PtV, formed while ~Pt5Al3 was present both in the as-cast and annealed conditions. The hardness of Alloy

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2H reduced by approximately 100 HV0.3, suggesting that one or both of the two phases that disappeared contributed to the higher hardness in the as-cast condition. The hardness of Alloys 3 and 4 remained more or less the same after annealing. In Alloy 3, the two phases ~Pt3V and ~Pt2Al disappeared and two new phases, ~Pt2V and ~PtV, formed while in Alloy 4 there was no change in the phases. In Alloy 5, a new phase, ~Pt2V, formed during annealing, while the two phases (Pt) and ~Pt3Al were retained. There was a slight increase of hardness from 420Âą20 HV0.3 to 482Âą14 HV0.3, suggesting that the new phase was responsible for the increase, as expected. The phases in Alloy 6 remained the same and, as expected, the hardness was also similar after annealing. The hardness of Alloys 7 and 8 remained the same after annealing if the standard deviations are taken into account. This is not surprising as the phases ~PtV, ~PtV3 and ~Pt5Al3 remained the same before and after annealing. The hardness of Alloy 9 decreased by approximately 100 HV, while the two phases (Pt) and ~Pt3Al remained the same in the as-cast and annealed conditions. The decrease in hardness may be explained by coarsening of the phases during annealing. Alloy 10 changed from two-phase ~Pt3Al + ~Pt2Al to single phase ~Pt3Al, while the hardness increased from 462Âą14 HV0.3 to 535Âą23 HV0.3. This suggests that ~Pt2Al is a soft phase compared to ~Pt3Al and its disappearance contributed to the increase in hardness. Alloy 11 consisted of a single ternary phase, Ď„2 and was the hardest alloy, even in the as-cast condition, with Vickers hardness of 886Âą23 HV0.3. Further evidence of its hardness and brittleness were the visible cracks around the indentation in the as-cast condition. It disintegrated while being removed from the resin mount and so it was not possible to measure the hardness in the annealed condition. Figure 6 shows hardness values superimposed on isothermal section of Pt-Al-V at 1000°C, at the Pt-rich corner as derived and reported earlier (Odera et al., 2014). The same trend is seen as with the as-cast alloys, where the hardness values reduced with increasing Pt content.

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A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys The alloys having ternary phase dendrites (Alloys 17, 18 and 21) had very high hardness values, as expected. Alloy 18 had the highest hardness value (1109Âą25 HV0.3). The second hardest alloy was Alloy 21, which had a hardness of 1086Âą28 HV0.3. All the indentations on Alloy 15 were on the eutectic areas and the high hardness value of 936Âą59 HV0.3 reflects the high hardness of the ternary phase component of the eutectic, rather than the binary intermetallic component. The single-phase Alloys 12 and 13, which had high Pt contents, had the lowest hardness values. Alloy 16, which also had high Pt content, had a similarly low hardness. The rest of the alloys all had relatively high hardness values, except for Alloy 24 which had a low hardness value of 358Âą19 HV0.3. The hardness values were comparable to those obtained by SĂźss (2007) in his investigation of Pt-Al-Cr alloys. SĂźss (2007) used a load of 10 kg and many of the alloys cracked and deformed around the indentations. However, the highest Vickers hardness values obtained by SĂźss (2007) were between 800 HV10 and 900 HV10. The hardness of the Pt-CrV alloys investigated in this work were also generally much higher than the ternary Pt-Al based alloys investigated by Hill et al. (2001) which had a maximum hardness value of 530 HV (the load used was not specified). These hardness values for the alloys, whether in the as-cast or annealed condition, were much higher than for pure platinum, which is very soft and ductile with a Vickers hardness of approximately 50 HV (Murakami ,2008). Figure 7 shows hardness values superimposed on the solidification projection of the Pt-Cr-V system. It shows that alloys having high Pt content have the lowest hardness and there is an increase in hardness as one moves away from the Pt-rich corner.

There was a general increase in hardness values after annealing the Pt-Cr-V alloys, except for Alloys 15H and 23H, which had lower values and this was probably due to microstructure coarsening. Alloys 17, 18 and 21 (as-cast) had ternary phase dendrites with high hardness values in the

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as-cast condition, as expected. These three alloys also exhibited extreme brittleness. The annealed Alloys 18H and 21H cracked and disintegrated while being removed from the resin mounting. Alloy 18H cracked, but stayed as one piece and therefore it was possible to undertake microhardness tests. The annealed Alloy 20H also cracked and disintegrated while being removed from the resin mounting. The alloys with high Pt content had the lowest hardness values. If the three annealed alloys that disintegrated are included among those with hardness values exceeding 600 HV0.3, then approximately 62% of the alloys had Vickers hardnesses higher than 600 HV0.3. This shows that the alloys were generally harder than those investigated by SĂźss (2007), just over half of which had Vickers hardnesses higher than 600 HV10. Again, if the same three alloys are included among those that had hardness values higher than 750 HV0.3, then slightly more than half the alloys had hardness values more than 750 HV0.3 compared to those investigated by SĂźss (2007), only three of which (17%) had hardness values higher than 750 HV10. The assumption that the three alloys that disintegrated had hardness values higher than 750 HV0.3 is valid because the hardest annealed alloy whose hardness was measured is Alloy 18H (1086 HV0.3) and an other three alloys had hardness values exceeding 800 HV0.3. These alloys were much harder than the Pt-Al based alloys investigated by Hill (2001), the hardest of which had a Vickers hardness of 530 HV (the load used was not specified). The hardness of Alloys 12H and 13H remained similar after annealing. Alloy 12H remained single phase ~Pt3V after annealing, while Alloy 13H changed from a single phase structure of ~Pt3V to a two-phase structure of ~Pt2V and ~Pt3V. This shows that the two phases ~Pt2V and Pt3V have similar hardness. Alloy 16H, which was also single phase ~Pt3V in the as-cast condition, retained the same phase after annealing, with a slight increase in hardness. This is probably due to traces of (Pt) detected by XRD in the as-cast alloy but which disappeared after annealing. The phases in Alloy 14H were the same before and after annealing and there was not much change in the hardness, as expected. However, there was a substantial decrease in the hardness of Alloy 15H after annealing, although the phases were the same. This decrease is attributed to coarsening. In Alloy 17H, the phase ~Cr3Pt disappeared and ~Pt2V formed after annealing, while the ternary phase, Ď„3, was present both in the as-cast and annealed conditions. The increase in hardness from 854Âą30 HV0.3 to 1085Âą46 HV0.3 is therefore attributed to ~Pt2V, as expected. Alloy 18, which had the highest hardness value at 1109Âą25 HV0.3 in the ascast condition, disintegrated while being removed from the resin mounting after SEM analysis following annealing, therefore its hardness was not measured. However, the phases, Ď„3, A15(~Cr3Pt) and ~Cr3Pt, remained the same in the as-cast and annealed conditions and the disintegration indicated extreme hardness and brittleness of the alloy. The phases in Alloy 19H did not change after annealing, but the hardness increased from 796Âą35 HV0.3 to 923Âą 57 HV0.3. This was the same with Alloy 20H, where the phases (V,Cr) and ~PtV3 remained the same in the as-cast and annealed conditions. However, Alloy 20H disintegrated while being removed from the resin mounting and it was not VOLUME 117

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A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys possible to measure its hardness after annealing, although the disintegration suggested extreme hardness and brittleness. Alloy 21 had a hardness value of 1086Âą28 HV0.3, making it the second hardest alloy in the as-cast condition. It had the same phases ~Cr3Pt, ~PtV3, Ď„3 and (V,Cr) in the as-cast and annealed conditions. It disintegrated while being removed from the resin mounting and as such, its annealed hardness was not measured. The disintegration indicated extreme hardness and brittleness. The phases, (V,Cr) and A15(~Cr3Pt), in Alloy 22H remained the same in the as-cast and annealed conditions, but the hardness increased from 707Âą41 HV0.3 to 837Âą46 HV0.3. Alloy 23H had the same phases, (V,Cr) and A15(~Cr3Pt) in the as-cast and annealed conditions, but the hardness decreased from 692Âą73 HV0.3 to 544Âą20 HV0.3. This deviated from the trend seen for most of the alloys in this system. Alloy 24H also retained the same phases, ~Cr3Pt and A15 (~Cr3Pt), in the as-cast and annealed conditions, but there was a substantial increase in hardness from 358Âą9 HV0.3 to 838Âą27 HV0.3. This was because heat treatment produced more of the A15 (~Cr3Pt), which would be expected to be harder, being an intermetallic phase. Figure 8 shows hardness values superposed on the isothermal section at 1000°C derived by and reported in Odera (2013). Hardness is lowest at the Pt-rich corner and increases with reduction in Pt content. Figure 9 shows a combined plot of hardness values of ascast and annealed Pt-Al-V alloys, while Figure 10 shows a similar plot for Pt-Cr-V alloys. There is a general increase in hardness values for the Pt-Cr-V alloys after annealing, while there is little change in the values for Pt-Al-V alloys.

01)/ 3,013 Microhardness values for selected Pt-Al-V and Pt-Cr-V alloys were determined in the as-cast and annealed conditions. Alloys with high Pt content (more than 70 at.%) had the lowest hardness values, while those containing ternary phases had the highest hardness. The single ternary phase Alloy 11 (average composition Pt56.9:Al25.4:V17.7) had the highest hardness of 886Âą23 HV0.3 among the Pt-Al-V alloys in the as-cast condition. Alloy 18 (average composition Pt25.8:Cr42.9:V31.3) had the highest hardness value of 1109Âą25 HV0.3 among the Pt-Cr-V alloys in the as-cast condition. These two sets of alloys, especially the Pt-Cr-V alloys, had higher hardness values than the Pt-Al-Cr alloys studied by SĂźss (2007) and Pt-Al based alloys studied by Hill et al. (2001), as well as the high-order Pt-based alloys studied by Odera et al. (2015).

;2 2-21)23 CORNISH, L.A., FISCHER, B. and VÖLKL, R. 2003. Development of platinum-groupmetal superalloys for high-temperature use. MRS Bulletin, vol. 28, no. 9. pp. 632–638. DIETER, G.E. 1986. Mechanical Metallurgy. 3rd edn. McGraw-Hill, New York. pp. 329–330. HILL, P.J., CORNISH, L.A., ELLIS, P. and WITCOMB, M.J. 2001a. The effects of Ti and Cr additions on the phase equilibria and properties of (Pt)/Pt3Al alloys. Journal of Alloys and Compounds, vol. 322, no. 1. pp.166–175. HILL, P.J., BIGGS, T., ELLIS, P., HOHLS, J., TAYLOR S.S. and WOLFF, I.M. 2001b. An assessment of ternary precipitation-strengthened Pt alloys for ultra-high temperature applications. Materials Science and Engeering A, 301, pp. 67–179. MURAKAMI, T., SAHARA, R., HARAKO, D., AKIBA, M., NARUSHIMA, T. and OUCHI, C. 2008. The effect of solute elements on hardness and grain size in platinum based binary alloys. Materials Transactions, vol. 49, no. 3. pp. 538–547.

Table III shows the HV0.3 hardness values of higher order Ptbased alloys (Odera et al., 2015). Pt-Al-V and Pt-Cr-V alloys were generally much harder than the as-cast higher order alloys, because many of the Pt-Al-V and Pt-Cr-V alloys contained hard Pt-V intermetallic phases. The high hardness of some of the Pt-Al-Cr alloys (SĂźss, 2007) was attributed to ~PtAl2, ~PtAl and ~Pt2Al3.

, -25 0 ,12+5'4-+12335#/0.50 543()43.541+541124/2+5 .(!/(& 4//0"3

, -25 4-+12335 4/ 235$ & * %50 541124/2+5 .( -(&54//0"3 3 #2-#032+5015.'25,30.'2- 4/532).,0154.5 5$4.* %

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, -25 0 ,12+5'4-+12335#/0.50 543()43.541+541124/2+5 .( -(& 4//0"3

A comparative study of hardness of Pt-Cr-V and Pt-Al-V alloys Table III

4-+12335 4/ 2350 543()43.541+541124/2+5', '2-50-+2-5$ +2-45 * 5 %54//0"3 !//0"5 5 0*

! 2-4 25 )0 #03,.,015$4.* %

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'43235#-2321.5,15 .'2543()43.54//0"3

'43235#-2321.5 ,15.'2541124/2+54//0"3

1 and 1H

Pt63.9:Al12.2:Cr4.3:Ru0.7:V18.9

537Âą13

700Âą20

~Pt3Al, (Pt)

~Pt3Al, ~Pt2V

2 and 2H

Pt69.5:Al11.5:Cr4.2:Ru0.6:V14.2

428Âą11

482Âą24

~Pt3Al, (Pt)

~Pt3Al

3 and 3H

Pt75.2:Al11.2:Cr4.0:Ru0.6:V9.5

377Âą8

359Âą9

~Pt3Al

~Pt3Al

4 and 4H

Pt78.7:Al12.2:Cr3.8:Ru0.6:V5.2

422Âą9

-

~Pt3Al

~Pt3Al, (Pt)

5 and 5H

Pt63.2:Al12.9:Cr4.0:Ru0.7:V19.0:Nb0.6

603Âą21

821Âą32

~Pt3Al, (Pt)

~Pt3Al, ~PtV, ~Pt2V

6 and 6H

Pt71.7:Al12.8:Cr4.9:Ru1.1:V9.9:Nb0.3

545Âą12

535Âą15

~Pt3Al, (Pt)

~Pt3Al

ODERA, B.O. 2013. Addition of vanadium and niobium to platinum-based alloys. PhD thesis, University of the Witwatersrand, Johannesburg. ODERA, B.O., CORNISH, L.A., PAPO, M.J. and RADING, G.O. 2012a. Electrolytic etching of platinum-aluminium based alloys. Platinum Metals Review, vol. 56, no. 4. pp. 257–261. ODERA, B., CORNISH, L.A., SHONGWE, M.B., RADING, G. and PAPO, M.J. 2012b. As cast and heat-treated alloys of the Pt-Al-V system at the Pt-rich corner. Journal of the Southern African Institute of Mining and Metallurgy, vol. 112, no. 7. pp. 505–516. ODERA, B., PAPO, M.J., RADING, G. and CORNISH, L.A. 2014. Experimental solidification projection, liquidus surface projection and isothermal section at 1000° C for the Pt-Cr-V system. Journal of Phase Equilibria and Diffusion, vol. 35, no. 4. pp. 476–489. ODERA, B., PAPO, M.J., COUPERTHWAITE, R., RADING, G., BILLING, D. and CORNISH, L.A. 2015. High-order additions to platinum-based alloys for

high-temperature applications. Journal of the Southern African Institute of Mining and Metallurgy, vol. 115, no. 3. pp. 241–250. SIMS, C.T., STOLOFF, N.S. and HAGEL, W.C. 1987. High-temperature Materials for Aerospace and Industrial Power. Wiley, New York. SĂźss, R., Cornish, L.A., Hill, P.J., Hohls, J. and Compton, D.N. (2003). Properties of a new series of superalloys based on Pt80:Al14:Cr3:Ru3. Advanced Materials and Processes for Gas Turbines, pp. 301–307. SĂœSS, R. 2007. Investigation of the Pt-Al-Cr system as part of the development of the Pt-Al-Cr-Ru thermodynamic database. PhD thesis, University of the Witwatersrand, Johannesburg. SĂœSS, R., WATSON, A., CORNISH, L.A. and COMPTON, D.N. 2009. Development of a database for the prediction of phases in Pt–Al–Cr–Ru alloys for hightemperature and corrosive environments: Al–Cr–Ru. Journal of Alloys and Compounds, vol. 476, no. 1. pp. 176–186. N

How it works:

5 Star Incentive Programme

1. Collect your membership card from the SAIMM. 2. Visit the SAIMM Website partner page to view all the deals available 3. Present it to the selected providers on the page behind this to qualify for the discount or offer.

About the SAIMM 5 Star Incentive Programme: The SAIMM are proud to welcome you to our Incentive programme where we have negotiated to provide you with more benefits.

Our Partners

These benefits include: 1. Top 5 Proposers for the current financial year are to be given a free ticket to the SAIMM Annual Banquet with mention at the Annual General Meeting. 2. Top 5 Referees for the financial year are to be given a free ticket to the SAIMM Annual Banquet with mention at the Annual General Meeting. 3. Access to discounts offered by Service providers that have negotiated discounted rates and special offers for you our valued member. 4. Conference attendance within a 2 year period and applies to events that are paid for. If you attend 3 events you get the next conference that you register for at no cost.

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5. The author with the most number of papers published in the SAIMM Journal in the previous financial year would be recognised at the Annual General Meeting and will receive a free ticket to the SAIMM Annual Banquet.

11 June 2018 — Workshops • 12–13 June 2018 — Conference 14 June 2018 — Technical Visits

Birchwood Conference Centre (Jet Park, Johannesburg) BACKGROUND eing the seventh conference in the series, the Diamonds – Source to Use conferences target the full spectrum of the diamond pipeline from exploration through to sales and marketing. The last conference in this series (Diamonds – Still Sparkling) was held in Botswana in 2016 and it is now returning to Johannesburg, where it was last held in 2013

B

OBJECTIVE The objective of the conference is to provide a forum for the dissemination of information relating to the latest tools and techniques applicable to all stages of the diamond industry; from exploration, through mine design, processing, to cutting, marketing and sales. WHO SHOULD ATTEND Geologists Mineral (Diamond) Resource Managers Mining Engineers Process Engineers Consultants Suppliers Sales/marketing Diamantaires Mine Managers Mining Companies Students.

TOPICS Geology and exploration Mine expansion projects Mining, metallurgy and processing technology Rough diamond sales and marketing Cutting, polishing and retail Financial services and industry analysis Industry governance, beneficiation and legislation Mine-specific case-studies Value optimization.

For further information contact: Camielah Jardine • Head of Conferencing ¡ SAIMM P O Box 61127 • Marshalltown 2107 Tel: (011) 834-1273/7 Fax: (011) 833-8156 or (011) 838-5923 E-mail: camielah@saimm.co.za • Website: http://www.saimm.co.za

http://dx.doi.org/10.17159/2411-9717/2017/v117n10a10

Alternative carbon materials as practical and more durable fuel cell electrocatalyst supports than conventional carbon blacks by M.L. Stevenson and G. Pattrick

Conventional low-temperature proton exchange membrane fuel cell (LTPEMFC) electrocatalyst supports suffer from excessively high corrosion rates. In this work, several alternative carbons were investigated in terms of their viability to produce ‘drop-in’ alternative electrocatalysts to the conventional materials while providing significantly enhanced support corrosion resistance. This was to establish their potential to serve both current and future markets. The materials examined were: partially graphitised Vulcan, a hierarchical porous graphene-like material and commercial graphitised carbon nanofiberes (CNFs). Samples were prepared at 40 wt% Pt loading and performance compared to conventional Pt (40wt%)/Vulcan and Pt (40wt%)/Ketjenblack. CNFs were the only material to display improved corrosion resistance over the conventional supports. The catalyst did, however, suffer from comparatively low beginning-of-life performance. This work did nonetheless demonstrate good promise for the commercial CNF support in producing functional electrocatalysts that could (a) potentially be directly substituted into well-established conventional catalyst-based MEA manufacturing processes and (b) offer considerably enhanced support corrosion resistance. E,$CB;A PEM fuel cell, electrocatalyst support, corrosion, carbon nanfibres, durability.

H?FBC;:>F@C? Conventional electrocatalysts supported on Vulcan XC72R and Ketjenblack EC300J display poor durability under typical automotive operating conditions, as well as certain stationary applications (Zhang, 2008; Gasteiger, Vielstich and Lamm, 2003; Rabis, Rodriguez and Schmidt, 2012). This leads to fuel cell stack lifetimes far lower than current targets set by the US Department of Energy (5000 hours for automotive; 40 000 hours for stationary) (US Department of Energy, 2013). Two major degradation mechanisms, particularly on the cathode, are responsible for this: instability in the nanometre-sized Pt particles on the catalysts and corrosion of the carbon support material (Gasteiger, Vielstich and Lamm, 2003). Both of these mechanisms lead to losses in active area and, as a result, losses in catalytic activity of the materials. Corrosion of the support also reduces the porosity of the electrode catalyst layers, decreasing the accessibility of reagent gases to the active sites on the catalyst (Schulenburg et al., 2011). Typically, systems-level measures are incorporated into the fuel cell stack design to

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Ž High electronic conductivity, in excess of 1 S/cm. This is needed to keep electronic resistance (ohmic) losses in the catalyst layer to a minimum Ž High specific surface area, in the region of 50 to 800 m2/g, to achieve high dispersions of the active Pt metal. This is to maximise the utilisation of the catalyst Ž Ability to form a structure with high porosity when fabricated into catalyst layers. This is essential for creating a large three-phase interfacial area in the electrodes of an MEA. Increasing the graphitic content of a carbon material is a common technique to improve its stability. Yu et al., (2006) reported that with in situ testing, an MEA fabricated using a catalyst supported on a graphitised carbon black showed a degradation rate five times lower than that of an MEA made with a standard carbon-black-supported catalyst. This was after 1000 start/stop cycles – a degradation protocol particularly taxing to a carbon support.

* HySA/Catalysis, South African Council for Mineral Technology, Mintek, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2017. ISSN 2225-6253. This paper was first presented at the AMI Precious Metals 2017 Conference ‘The Precious Metals Development Network’ 17–20 October 2017, Protea Hotel Ranch Resort, Polokwane, South Africa.

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prevent modes of operation which are particularly damaging to the support (Zhang, 2008; Gasteiger, Vielstich and Lamm, 2003). Carbon corrosion does nevertheless occur and so to further improve the useful life of the membrane electrode assembly (MEA), catalysts supported on materials with inherently higher corrosion resistance are needed. Conventional supports possess several properties essential for high performance catalyst layers within MEAs (Rabis, Rodriguez and Schmidt, 2012), which alternative carbons would also need to maintain:

Alternative carbon materials as practical and more durable fuel cell electrocatalyst of selected catalyst samples. Co K radiation was used and scans were performed from 10° to 120° (2 ) using a Bruker D8 Advance diffractometer. Average Pt crystallite sizes were calculated using the Scherrer equation applied to the Pt(220) diffraction peak (Venkateswara Rao and Viswanathan, 2010).

High-resolution TEM images, using a JEOL JEM-2100F instrument, were taken of selected catalyst samples to visualise metal dispersions, for validation of the Pt surface areas predicted by cyclic voltammetry and also the average crystallite sizes calculated from XRD. .@3:BEG G E2F G)/*% G@=D3EGC2GDA+BE>E@ E;G6,BC3BD2G6/+#&+ *+))* >DB!C?G?D?C2@!BEA'G/@30F G8C=<:FEB+3E?EBDFE;GBE<BEAE?FDF@C?GC2 AFD> E;+>:< GAFB:>F:BEGC2GF0EG?D?C2@!BEAG5AC:B>E G0FF< --<,BC3BD2+ <BC;:>FA'>C= GD>>EAAE;G G E>E=!EBG#7 4

Highly graphitised carbon materials other than carbon black are also being investigated as viable alternative supports. In particular, graphitised carbon nanofibres and nanotubes have shown promise in producing catalysts with high metal dispersions and improved support corrosion resistance over the standard carbon blacks (Guha, Zawodzinski and Schiraldi, 2008; Lee et al., 2006; ZaragozaMartin et al., 2007). Another key consideration of this work, however, was to establish whether certain alternative supports would be capable of providing enhanced support corrosion benefits for both current and future applications. Since many existing large-scale MEA manufacturing processes are currently based on conventional catalysts, viable alternative carbons would need to produce electrocatalysts directly compatible with these processes to see rapid entry into current commercial pipelines. The alternative carbons examined were: a partially graphitised form of Vulcan XC72R (pG-Vulcan) produced by CabotÂŽ; highly graphitised carbon nanofibres (CNFs) from PyrographÂŽ Products, Inc. (PR-24-XT-HHT); and a 3D network, hierarchical porous graphene-like material (HPG) obtained through international collaborations within the HySA/Catalysis programme.

% <EB@=E?FD9 ! ! ! ! Catalysts were prepared according to a proprietary synthesis technique developed within the HySA/Catalysis programme. In brief, the routine involved the precipitation of Pt complexes from an aqueous solution of Pt precursor, followed by deposition on to the carbon support material and in situ reduction with a suitable reducing agent. Syntheses in per batch quantities from 0.5 g to 1 kg per batch have been thoroughly investigated and shown to yield Pt/C materials supported on Vulcan XC72R and Ketjenblack EC-300J without any losses in performance throughout the scale-up procedure.

! ! ! ! ! XRD was used to characterise the crystallographic structure

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BET surface area measurements were performed on the various support materials using a Micromeritics ASAP 2020. All samples were degassed at 350°C under vacuum for 4 hours prior to measurement.

All TF-RDE measurements were performed in 0.1M HClO4 using a three-electrode cell. A saturated calomel electrode (SCE) served as reference and a Pt strip as counter-electrode. All reported potentials were referenced back to the relative hydrogen electrode (RHE). The working electrode was a 5 mm glassy carbon RDE and all measurements were carried out on an Autolab PGSTAT302N potentiostat. Prior to the measurements, the glassy carbon electrodes were polished with 0.05 Îźm alumina on a Buehler TexmetÂŽ polishing pad. Catalyst inks were prepared by dispersing a calculated mass of catalyst in approximately. 15 mâ„“ of highpurity water by ultrasonication for 30 minutes. Catalyst was then deposited at a Pt loading density of 56 Âľg m2 after which 27.8 ÂľL of a 0.2817 mg/mL NafionÂŽ solution was micropipetted on top of the catalyst layer and dried in an oven at 120°C for 10 minutes. Pt surface area measurements were performed in N2purged, 0.1M HClO4, maintained at 25°C. The catalyst layers on the RDEs were first conditioned by cycling the potential between 0 and 1.2 VRHE at 50 mV/s for 15 cycles. Thereafter, three further cycles were measured at 20 mV/s and the hydrogen adsorption charge of the third cycle used to calculate the Pt surface area of the catalyst. In order to limit contributions from molecular hydrogen evolution, integration for the hydrogen adsorption charge was performed up until the potential of the current maximum following the Pt(111) adsorption peak in the cathodic sweep. A specific charge transfer value of 210 ÂľC/cm2 Pt was assumed. Kinetic ORR activity measurements were carried out in O2-saturated electrolyte at 60°C. The RDE was rotated at 1600 r/min and the potential was swept from 0 to 1.0 VRHE at 20 mV/s. Background correction to remove the capacitive currents not related to ORR was performed in N2-purged electrolyte at 60°C with zero rotation of the RDE. Mass transport limitations were removed using the KouteckyLevich equation (Gasteiger et al., 2005) and a Tafel analysis performed on the kinetic data between 875 and 950 m VRHE. Pt particle stability was assessed by cycling the potential between 0.6 and 1.0 VRHE for 7200 square-wave cycles, with Pt SA measured at 1200-cycle intervals. This was done in N2purged 0.1M HClO4 at 60°C. Support corrosion resistance was evaluated by holding the potential on the RDE at 1.5 VRHE for

Alternative carbon materials as practical and more durable fuel cell electrocatalyst 2 hours and then averaging the final 2 minutes of the measurement to obtain a measure of the support corrosion current. This was performed in N2-purged, 25°C HClO4.

Following this, the MEA was assessed by a number of standard in situ testing protocols, the conditions for which are also contained in Table I.

! !

/EA:9FAGD?;G;@A>:AA@C?

! !

Catalyst-coated substrates (CCS) with active areas of 50 cm2 were fabricated by a procedure developed within HySA/Catalysis. Catalyst inks were prepared using a proprietary combination of solvents, catalyst and 15 wt% ionomer solution (IonPower™, LQ-1015). An I:C ratio of 0.75 was employed for both the anode and cathode catalyst layers. A SilversonÂŽ L5MA high-sheer mixer was used to homogenise the ink components before transfer to a USIÂŽ Prism 300 spray coater. Catalyst layers were then deposited onto Freudenberg H2315 I2C6 GDLs by ultrasonic spray coating and dried in an oven at 120°C for 30 minutes. Targeted electrode Pt loadings were 0.4/0.4 mg Pt/cm2 (anode/cathode). The resulting gas diffusion electrodes (GDEs) were hot-pressed onto opposite sides of NafionÂŽ NRE211 membrane at a pressure of 10 MPa relative to active area, at 135°C for 3 minutes.

Figure 2 displays a comparison of the BoL thin-film RDE cyclic voltammograms measured on the catalysts synthesised with the various supports. Included are the CVs obtained on the conventional Pt(approx. 40 wt%)/Vulcan XC72R (HySAV40) and Pt(approx. 40 wt%)/Ketjenblack EC-300J (HySAK40) materials. Table II summarises the Pt loadings, BoL Pt surface areas and kinetic ORR activities measured on the various catalyst samples and Table III provides a comparison of the Pt particle sizes determined by TF-RDE, XRD and HRTEM. Also shown in Table III are the BET surface areas measured on the supports. What is immediately apparent in Figure 2 is the extremely low initial metal area measured on the CNF-supported catalyst relative to the other Pt (approx. 40 wt%) materials. At just 35.8 m2/g Pt, this represents only 41% of the metal area obtained on the standard Vulcan-supported HySA-V40. The TF-RDE Pt surface area ranking on the catalysts, as well as the XRD Pt crystallite and HRTEM Pt particle sizes in general, followed the BET surface areas measured on the supports, with lower BET SAs yielding lower Pt SAs. This is unsurprising, as smaller support surface areas relate to a lower number of active sites on the support for metal deposition, leading to a lower Pt dispersions for a given synthesis procedure.

Prior to evaluation, newly fabricated MEAs were subjected to a conditioning protocol in order to stabilize performance at the assembly’s peak value. This involved cycling the potential on the MEA between 1.0 V (or open circuit) for 30 seconds and 0.3 V for 10 minutes. This was performed for a total of 12 cycles under the conditions summarised in Table I.

Table I

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A/C Feed Gas A/C Feed Flow Rates A/C Back Pressure

8C?;@F@C?@?3

6C9DB@ADF@C?G8:B E8*/G G 77G=(->=#

@?EF@>G // (>F@ @F,

8DF0C;EG%61(),;BC3E?G8BCAAC EB Hydrogen/Nitrogen

Hydrogen/Air

Hydrogen/Air

Hydrogen/Oxygen

2.0 x Stoich. at 1600 mA/cm2

1.5/2.0

2.0/4.0 NL/min

0.5/0.5 NL/min

2.0/2.0 bar(abs)

2.0/2.0 bar(abs)

1.0/1.0 bar(abs)

1.0/1.0 bar(abs)

A/C Feed Gas Temp. A/C Feed Humidity

90/90 °C

90/90 °C

90/90 °C

90/90 °C

100/80% RH

100/80% RH

100/100 %RH

100/100 %RH

80°C

80°C

80°C

80°C

Cell Temp.

Table II

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HySA-V40 (FCU4007) HySA-K40 (FCKU4007) FCHPG4002

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6F

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Vulcan XC72R

50

37.58 Âą0.08

87.4 Âą0.9

66.4 Âą0.8

220 Âą5

0.192 Âą0.005

Ketjenblack

50

42.18 Âą0.14

110.2 Âą4.7

62.8 Âą1.3

152 Âą9

0.168 Âą0.006

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1<E>@2@> 5 (->=# 6F4

DAA 5(-=3G6F4

HPG

7

40.02

97.9 Âą6.1

70.3 Âą2.1

189 Âą12

0.186 Âą0.022

FCGV4001

pG-Vulcan

7

38.78 Âą0.14

75.5 Âą5.3

66.2 Âą2.9

258 Âą47

0.193 Âą0.023

FCNF4003

CNFs

7

38.30 Âą0.12

35.8 Âą0.5

64.7 Âą0.6

324 Âą20

0.116 Âą0.008

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Alternative carbon materials as practical and more durable fuel cell electrocatalyst Table III

8C=<DB@A@C?GC2G6FG<DBF@>9E->B,AFD99@FEGA@ EAG;EFEB=@?E;G!,G*.+/ % G / GD?;G)/*% C?GF0EG DB@C:AG>DFD9,AFGAD=<9EA GD9C?3G$@F0GF0EG %*GA:B2D>EGDBEDAG=EDA:BE;GC?GF0E A:<<CBFA 1D=<9EG D=E

1:<<CBF

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5=#-3G1:<<CBF4

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/ 5?=4

)/*% 5?=4

Vulcan XC72R

237.1

3.2

3.3

2-4

Ketjenblack

850.6

2.5

2.7

2-3

HPG

581.3

2.9

3.6

2-4

FCGV4001

pG-Vulcan

147.9

3.7

-

2-4

FCNF4003

CNFs

44.7

7.8

4.6

3-5

HySA-V40 (FCU4007) HySA-K40 (FCKU4007) FCHPG4002

The average metal particle size calculated from the TFRDE Pt SA on the Pt/CNF catalyst was 7.8 nm. HRTEM imaging of the catalyst, however, displayed a high dispersion of the Pt metal across the nanofibres with Pt particle sizes typically in the 3 to 5 nm range (Figure 3). In agreement with this is the average Pt(220) crystallite size estimated from XRD, which was 4.6 nm (Figure 4). The reason for the large disparity between these measurements on the CNF-supported catalyst could be the possibility of Pt particle growth inside the nanofibre tubes during synthesis, which may have been inaccessible to the electrolyte during RDE measurements.

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HRTEM images appear to confirm this, showing a clear line of Pt metal particles along the inner wall of the CNFs (Figure 3). In terms of kinetic ORR activity measurements, specific activity was approximately 47% higher on the Pt(38.30 wt%)/CNF sample than the HySA-V40, but because of the very low Pt SA, mass activity was only 60% of the standard Vulcan-supported catalyst at 0.116 A/mg Pt. Moving on to the Pt(40.02wt%)/HPG catalyst, as can be seen in Table II, it possessed an initial Pt surface area between that of the HySA-V40 and HySA-K40 catalysts at 97.9 m2/g Pt. In agreement with this, HRTEM imaging of the catalyst indicated particles in the 2-4 nm range. Kinetic ORR mass activity was found to be essentially the same as the HySA-V40. BoL performance of the HPG supported catalyst therefore seemed very promising, matching that of the conventional catalysts. The pG-Vulcan-supported catalyst showed a Pt surface area only marginally lower than the conventional HySA-V40. Kinetic ORR mass activity was also on par with the HySAV40 with, however, a slightly higher specific activity at 900 m VRHE, attributed to the larger average Pt particle size on the pG-Vulcan-supported material. This was possibly due to the Pt particle size effect. Being referred to here is the straightforward relationship between Pt particle size and ORR activity, where larger particles have been associated with higher specific activities and lower mass activities and smaller particles associated with smaller specific activities and higher mass activities

Alternative carbon materials as practical and more durable fuel cell electrocatalyst

(Gasteiger et al., 2005). The drop-off in mass activity with increasing Pt particle size has been found to increase somewhat only at values below approximately 60-70 m2/g Pt (Gasteiger et al., 2005). Thus with sample FCGV4001 possessing approximately 75.5 m2/g Pt, the impact on mass activity of its slightly larger particles on average than the HySA-V40 was perhaps not seen because of this phenomenon. Looking now at durability testing, Figure 6 presents the Pt surface area decay profiles obtained from the thin-film RDE Pt stability ADT. As can be seen, with the exception of the HPG-supported catalyst, all the materials actually showed fairly similar degradation profiles, losing between 35 and 42% of their initial Pt surface area by the end of the test. The HPG-supported catalyst clearly showed the highest degradation, however, losing nearly 60% of its active area by the end of the ADT. The CNF- and pG-Vulcan supported catalysts performed only as good as their standard Vulcan-and Ketjenblacksupported counterparts and so improved support durability would be the only potential advantage the alternative carbons could offer over the conventional carbon blacks. Figure 7 displays the corrosion currents measured over the final 1200 seconds of the TF-RDE support corrosion resistance ADT. In terms of the alternative carbons, the HPG-supported catalyst showed the highest corrosion current, which was in fact 10% higher than the Pt(approx. 40 wt%)/Ketjenblack material. Showing both lower Pt particle stability and support corrosion resistance, the HPG material did not appear to have any improved durability over the conventional carbon blacks and therefore would likely not provide any significant advantage over these materials.

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The partially graphitised Vulcan-supported catalyst, rather surprisingly, showed essentially the same average corrosion current as the HySA-V40. It was anticipated that the graphitisation imparted into the material, however small, would at least provide a noticeable enhancement in support corrosion resistance. With respect to the TF-RDE BoL and ADT measurements, therefore, the pG-Vulcan material was almost indistinguishable from standard Vulcan, except in terms of BoL Pt SA, for which it showed a 14% reduction. It was thus expected that it would show no appreciable advantage over the HySA-V40 in a fuel cell environment either and with the material costing more than standard Vulcan, it would in fact be less attractive as a fuel cell catalyst support.

! ! ! ! ! ! The CNF-supported catalyst was the only material to show any improvement in support corrosion resistance over the conventional carbon blacks. The improvement was in fact fairly substantial, with the two-minute average corrosion current approximately 36 and 63% lower than the HySA-V40 and HySA-K40 catalysts, respectively. Unfortunately, this improved support corrosion resistance came at the cost of a large reduction in BoL Pt dispersion and kinetic ORR activity. In view of these results, it was decided to investigate lowering the Pt loading on the CNF supported catalyst to 2 0wt%. Table IV summarises a comparison between the BoL TF-RDE measurements performed on the 20 and 40 wt% Pt/CNF catalysts, as well as the HySA-V20 and V40 materials. As can be seen, the 20 wt% Pt/CNF catalyst did indeed show an increase in Pt SA by approximately 60% over the 40 wt% material, achieving just under 60 m2/g Pt versus the 35.8 m2/g Pt of sample FCNF4003. Rather surprisingly, kinetic ORR mass activity was actually somewhat lower, at around 0.095 A/g Pt, compared to 0.115 A/gPt for the 40 wt% catalyst. A near 50% reduction in specific activity appeared to be the cause for this and thus a phenomenon besides the Pt particle size effect was likely at play. As can be seen in the HRTEM images of Figure 8, there appeared yet again to be a significant deposition of Pt particles on the inner nanofibre walls of the Pt(approx. 20 wt%)/CNF catalyst. A possibility therefore exists that a larger fraction of the total metal may have been inactive on the 20 wt% catalyst than the 40 wt%, resulting in the apparent lower mass activity measured on the material. VOLUME 117

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Alternative carbon materials as practical and more durable fuel cell electrocatalyst Table IV

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Vulcan XC72R

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37.58 Âą0.08

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CNFs

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0.116 Âą0.008

FCNF2001

CNFs

14

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Shown in Figure 9 is a comparison of the H2/air polarisation curves measured on the Pt/CNF- and HySA-K40based MEAs. Immediately apparent is the drastically lower performance of the CNF-supported catalysts compared to the benchmark HySA-K40 catalyst across the entire range of current densities measured. At 0.6 V, both the 20 and 40 wt% Pt/CNF materials displayed current densities less than half that of the HySA-K40, achieving approximately 570 mA/cm2 versus 1188 mA/cm2 on the conventional catalyst. A number of possible causes for this are evident in Table V, where, in particular, the in situ kinetic ORR mass activities measured on samples FCNF2001 and FCNF4003 were only 13 and 16% of the value obtained on the HySAK40. What can also be seen is that the slope in the ohmic region of the polarisation curves for the CNF-based MEAs was higher than for the HySA-K40. Since the HFR values were similar for the MEAs across the range of current densities, it is likely that increased mass transport resistance in the Pt/CNF based layers was the cause for this, having a greater effect on polarisation performance at lower current densities than in the conventional catalyst. Validation for the increased mass transport resistance is provided by a comparison between the maximum current densities produced – the Pt/CNFs only achieving 1300 mA/cm2 vs. 1600mA/cm2 for the HySA-K40. A reason for this could have been that the catalyst layer structure produced by the nanofibres had a smaller porosity than the Ketjenblack catalyst, increasing the resistance to diffusion of the reagent gases to the active sites and leading to a lower overall three-phase interfacial area. The bulk density of the CNF material was also noted as being exceptionally lower than that of Ketjenblack and it is therefore likely that the electrode catalyst layers were also appreciably thicker, resulting in poorer gas transport through the layers.

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Also apparent was a lower OCV on the Pt/CNF CCSs. Although the exact reason for this is unclear, it may have simply been related to the poorer ORR kinetics on the Pt/CNF cathode catalyst layers. In line with the poorer ORR kinetics was the charge transfer resistance at 100 mA/cm2, which was also higher on the CNF-based MEAs and cathodic EPSAs were 31 and 65% below that of the HySA-K40 for the 20 and 40 wt% CNF catalysts, respectively.

8C?>9:A@C?A Several alternative carbon materials were investigated in this work for their viability as replacement supports with improved durability over the conventional Vulcan XC72R and Ketjenblack EC-300J materials. BoL thin-film RDE measurements on samples synthesised at the 40 wt% Pt loading revealed that the HPG and pG-Vulcan catalysts had similar kinetic ORR mass activity at 900 mVRHE as the benchmark conventional catalysts. The Pt(approx. 40 wt%)/CNF catalyst, however, presented only 41% and 60% respectively of the Pt SA and kinetic ORR activity of the Pt(approx. 40 wt%)/Vulcan material. TF-RDE support corrosion resistance testing indicated that the HPG-supported catalyst had a corrosion current higher than both the V40 and K40 catalysts. Pt stability was also drastically lower, ruling out the material as a valuable, more durable replacement for the conventional carbon blacks. The pG-Vulcan catalyst presented essentially the same corrosion current as the V40; however, the CNF support showed a substantial decrease over the conventional materials, with a corrosion current approximately 63% below the HySA-K40.

Alternative carbon materials as practical and more durable fuel cell electrocatalyst Table V

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FCNF2001 (Pt(18.71wt%)/CNF) 0.4/0.4 NRE211 Freud. C6

1188 756 92 530

570 490 108 693

576 451 90 893

76 0.137 477 29 2.82

66 0.018 181 10 2.48

62 0.022 110 20 6.67

H2/air Pol. Curve - Current Density @ 0.6V (mA/cm2) H2/air Pol. Curve - Peak Power Density (mW/cm2) H2/air Pol. Curve - HFR @ 1000mA/cm2 (mΊ.cm2) Charge Transfer Resistance @ 100mA/cm2 (mΊ.cm2) ORR Kinetic Activity: Tafel Slope >850mV (mV/dec) Mass Activity @ 900mV (A/mg Pt) Specific Activity @ 900mV (ΟA/cm2 Pt) EPSA of Cathode (m2/g Pt) Hydrogen Crossover (mA/cm2)

To address the extremely poor BoL performance, the Pt loading on the CNF supported catalyst was reduced from 40 to 20 wt%. This, as expected, resulted in an increase in BoL Pt SA by nearly 60%, but with no corresponding improvement in kinetic ORR mass activity. This was proposed to be the result of Pt particles being deposited on the inner tube walls of the CNF material, with a larger fraction of the total Pt metal occurring on the inner walls of the Pt(20 wt%)/CNF sample. In situ H2/air polarisation curve performance was appreciably lower on both the 20 and 40 wt% Pt/CNF catalysts compared to the HySA-K40. The most significant reasons for this appeared to be in situ kinetic ORR mass activity, which was less than 20% of the value measured on the Ketjenblacksupported catalyst, as well as suspected higher mass transport resistance The CNF catalysts did, however, produce functional electrocatalyst layers showing modest performance without any modification of the MEA manufacturing procedure developed specifically for conventional Vulcan- and Ketjenblacksupported catalysts. There does therefore appear to be some promise for the CNFs in being able to produce electrocatalysts that could function as drop-in replacements for the conventional catalysts. In doing so, they could benefit not only future applications, but current applications as well. A deeper consideration of the surface properties of the supports and resulting catalysts (microscopic and macroscopic structures, effects on the interparticle distance between the Pt particles of the catalysts, etc.) may also help in further elucidating the various phenomena observed in terms of both the ex-situ and in situ measurement of the various supports. This was beyond the scope of this work, but may be addressed in a follow-up at a later stage.

(> ?C$9E;3E=E?FA The authors would like to thank the South African Department of Science and Technology (DST) for funding this research through the HySA/Catalysis programme. We would also like to thank Professor Pei Kang Shen from the Advanced Energy Materials Laboratory at the Sun Yat-Sen University of China for providing the hierarchical porous graphene sample material.

/E2EBE?>EA CARTER, R..N., KOCHA, S.S., WAGNER, F.T., FAY, M. and GASTEIGER, H.A. 2007. Artefacts in measuring electrode catalyst area of fuel cells through cyclic voltammetry. ECS Transactions, vol. 11, no. 1. pp. 403–410. GASTEIGER, H.A, VIELSTICH, W. and LAMM, A. 2003. Handbook of Fuel Cells – Volume 3: Fuel Cell Technology and Applications. Wiley, Chichester, UK. GASTEIGER, H.A., KOCHA, S.S., SOMPALLI, B. AND WAGNER, F.T. 2005. Activity benchmarks and requirements for Pt, Pt-alloy and non-Pt oxygen reduction catalysts for PEMFCs. Applied Catalysis B: Environmental, vol. 56, no. 1-2. pp. 9–35. GUHA, A., ZAWODZINSKI, T.A. and SCHIRALDI, D.A. 2008. Influence of carbon support microstructure on the polarization behavior of a polymer electrolyte membrane fuel cell membrane electrode assemblies. Journal of Power Sources, vol. 195, no. 16. pp. 5167–5175. LEE, K., ZHANG, J., WANG, H. and WILKINSON, D.P. 2006. Progress in the synthesis of carbon nanotube- and nanofiber-supported Pt electrocatalysts for PEM fuel cell catalysis. Journal of Applied Electrochemistry, 36 (5), pp. 507–522. LI, Y., LI, Z. and SHEN, P. 2013. Simultaneous formation of ultrahigh surface area and three-dimensional hierarchical porous graphene-like networks for fast and highly stable supercapacitors. Advanced Materials, vol. 25 no. 17. pp. 2474–2480. RABIS, A., RODRIGUEZ, P. and SCHMIDT, T. 2012. Electrocatalysis for polymer electrolyte fuel cells: recent achievements and future challenges. ACS Catalysis, vol. 2, no. 5. pp. 864-890. SCHULENBURG, A., SCHWANITZ, B., LINSE, N., SCHERER, G.G., WOKAUN, A. and KRBANJEVIC, J. 2011. Quantification of platinum deposition in polymer electrolyte fuel cell membranes. Electrochemistry Communications, vol. 13 no. 9. pp. 921–923. US DEPARTMENT OF ENERGY. 2013. Fuel cell technical team roadmap. http://energy.gov/sites/prod/files/ 2014/02/f8/ fctt_roadmap_june2013.pdf [accessed 15 December 2016]. VENKATESWARA RAO, C. and VISWANATHAN, B. 2010. Monodispersed platinum nanoparticle supported carbon electrodes for hydrogen oxidation and oxygen reduction in proton exchange membrane fuel cells. Journal of Physical Chemistry C, vol. 114, no. 18. pp. 8661–8667. YU, P.T., GU, W., MAKHARIA, R., WAGNER, F.T. and GASTEIGER, H.A. 2006. The impact of carbon stability on PEM fuel cell startup and shutdown voltage degradation. ECS Transactions, vol 3, no. 1. pp. 797–809. ZARAGOZA-MARTIN, F., SOPEÑA-ESCARIO, D., MORALLÒN, E. and SALINAS-MART�NEZ DE LECEA, C. 2007. Pt/Carbon nanofibers electrocatalysts for fuel cells: effect of the support oxidizing treatment. Journal of Power Sources, vol. 171, no. 2. pp. 302–309. ZHANG, J. 2008. PEM Fuel Cell Electrocatalysts and Catalyst Layers, SpringerVerlag, London, UK. N VOLUME 117

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Copper Cobalt Africa

In Association with

The 9TH Southern African Base Metals Conference

8–11 July 2018 Avani Victoria Falls Resort, Livingstone, Zambia

SAIMM is proud to host the Second Copper Cobalt Africa Conference in the heart of Africa To be held in Livingstone, Zambia, this anticipated and prestigious event provides a unique forum for discussion, sharing of experience and knowledge, and networking for all those interested in the processing of copper and cobalt in an African context, in one of the world’s most spectacular settings - the Victoria Falls. The African Copper Belt is again experiencing a resurgence of activity following the commodity downturn of recent years. A significant proportion of capital spending, project development, operational expansions, and metal value production in the Southern African mining industry take place in this region. The geology and mineralogy of the ores are significantly different to those in other major copper-producing regions of the world, often having very high grades as well as the presence of cobalt. Both mining and metallurgy present some unique challenges, not only in the technical arena, but also with respect to logistics and supply chain, human capital, community engagement, and legislative issues. This conference provides a platform for discussion of a range of topics, spanning the value chain from exploration, projects, through mining and processing, to recycling and secondary value addition. For international participants, this conference offers an ideal opportunity to gain in-depth knowledge of and exposure to the African copper and cobalt industries, and to better understand the various facets of mining and processing in this part of the world. Jointly hosted by the Mining and Metallurgy Technical Programme Committees of the Southern African Institute of Mining and Metallurgy (SAIMM), this conference aims to: Promote dialogue between the mining and metallurgical disciplines on common challenges facing the industry

Enhance understanding of new and existing technologies that can lead to safe and optimal resource utilization Encourage participation and build capacity amongst young and emerging professionals from the Copper Belt region. The organizing committee looks forward to your participation.

Conference Topics

Mineralogy Exploration and new projects Open-cast mining Underground mining Rock engineering Mine planning Mineral resource management Minerals processing Pyrometallurgy Hydrometallurgy Current operational practices and improvements Project development and execution Novel technologies for this industry Socio-economic challenges Recycling Waste treatment and minimization Environmental issues Geometallurgical modelling

To add your name to the mailing list and to receive further information, please e-mail your contact details to: Gugulethu Charlie Conference Co-ordinator Saimm, P O Box 61127, Marshalltown 2107, Tel: +27 (0) 11 834-1273/7 E-mail: gugu@saimm.co.za, Website: http://www.saimm.co.za

ASM Conference 2018 ‘A focus on Illegal, traditional, customary and Small Scale Mining in Sub-Saharan Africa’ 13–14 SEPTEMBER 2018 Birchwood Conference Centre, Boksburg BACKGROUND

C o n f e r e n c e

There are misinterpretations as regards Artisanal and Small Scale Mining (ASSM) often relating to a confusion and conflation of those engaging in ASSM and/or illegal mining. While there has been a lot of research in this field, often the researches overlap one another and it is site specific. There is a lack of clarity about legal versus illegal, and appropriate management practice. Often, mining engineers are placed in areas where there is ASSM but not have the necessary tools to deal effectively the sector. In addition, those engaging in ASSM do not know how to deal with practitioners of Large Scale Mining (LSM) with the result that the inherent conflict often resulting in violent altercations. In the case of illegal mining, there is no clarity from government on best practice for engineers and geoscientists resulting in frustration. South Africa has lagged behind developments in the rest of Africa with possible devastating consequences to its economy and people. It has failed to acknowledge the precepts of the African Mining Vision and the Millennium Development Goals, especially as regards small scale mining and poverty alleviation respectively. It has now become incumbent on mining practitioners to use their knowledge of mining to ensure that small scale miners mine safely, effectively and efficiently to ensure the future viability of the South African mining industry, while managing illegal mining effectively. It is therefore incumbent on all stakeholders to learn from the experiences of others in other parts of the world, as well as the socio-economic impacts of recognizing/not recognising ASSMs. Unfortunately, as with most current practices in South Africa, mining practitioners would have to ‘go it alone’ as government dithers. As such, this conference would seek to bridge the gap between LSM and ASSM, to consider the policy/legislative changes that would be required to develop the ASSM sector, implement the National Development Plan and African Mining Vision, to consider the implications of a future mining industry in South Africa where ASSM is rejected, on the one hand, and where ASSM is encouraged and properly regulated on the other and ASSM as a possible tool of poverty alleviation. As such, this conference seeks to go beyond the sharing of knowledge and encompass the transfer of skills and experiences as well.

THE ACT‘S MAIN OBJECTIVES ARE TO: Inform mining practitioners of the importance of Small Scale Miners Consider techniques and skills that miners would require to deal with Small Scale Miners

A n n o u n c e m e n t

Development of policy/legislative considerations to cater for ASSM versus illegal mining

Develop lessons that LSMs can teach practitioners of ASSM to mine effectively and efficiently

Considering of innovative machinery and techniques to ensure that ASSM

can be done effectively and safely Distinguish practitioners of ASSM from Illegal Miners Considering ways of implementing the provisions of the African Mining Vision Considering the provisions of the National Development Plan to develop the mining industry in the country Discuss the possibility of ASSM as a tool of poverty alleviation

WHO SHOULD ATTEND Community engagement practitioners Community liaison officials at the mines Closure practitioners Communities Affected by large and small scale mining NGOs Academics, Trainers and Educators Health and Safety Practitioners ASM/SSM specialists and advisors Government Officials and Regulators (DMR, DST, DTI; D Home Affairs) Poverty Alleviation Specialists Futurists and forward thinkers

TOPICS: ASSM vs ILLEGAL MINING There is a conflation between illegal mining and ASSM (customary, traditional or permitted right) often with one being used to refer to the other. At no stage can it be expected that mining companies interact with and accept responsibility for the activities related to illegal mining. However, due to the conflation of ASSM and illegal mining, it would appear that mining companies are expected to do just that. As a result, a proper definition of ASSM is required. It should be noted that, while a generic definition would not be possible, general characteristics can be identified. There is a need to evaluate best practice when confronted with illegal mining activities.

The Stakeholders There are many stakeholders that are involved in the ASSM debate, some may yet need to be identified. LSM cannot adopt a laid back attitude as they are directly and indirectly affected by ASSM. LSM companies need to acknowledge their impact on ASSM and the people engaged therein. At its most basic, there is the ‘Triple I’ of mining – Illegal, Informal and Institutional and all need to be properly defined so that a meaningful discussion can be had and the appropriateness of the term ‘ASSM’ be fully considered.

An African Perspective The African Mining Vision, and its Implementation Plan, provides an Afrocentric vision of ASSM. South Africa has yet to provide its vision for the implementing of the African Mining Vision. There is an opportunity to formalise the debates on ASSM into a unique African vision tinged with a South African perspective. It is also opportune to develop a coherent policy on ASSM.

Economic Development South Africa is seeing a shrinkage of its mining sector. Simultaneously, it is faced with deeper ore-bodies and depleting minerals. Still, mining is a major contributor to the economy and the development of skills. Recognition of ASSM would ensure that mining remains relevant as a contributor to the economy and skills and knowledge transfer and development. An evidencedbased study would be helpful in this regard. Its effectiveness as a poverty alleviation tool should also be considered.

Historical Research The Mineral and Energy Policy Centre (MEPC) had conducted major research into ASSM and how it can be a contributor to development in South Africa. This research has been neglected and is now outdated though very much relevant. A baseline study is required to determine the impact of ASSM, and its links, to the local economy.

Cultural Issues Mining is not an activity that commenced when colonisation occurred. There is evidence of mining prior to colonization. Studies would need to take account of these cultural practices as regards mining as it would impact the heritage of the people in the area, especially where LSM is taking place.

Beneficiation Where LSM is winding down, or where it is uneconomical to have LSM, ASSM should be an alternative to value addition of the minerals to be mined. This would impact the development of the area as well as creating opportunities for the beneficiation of products that would be relevant to the area and the community. There is minimal consideration of these issues at present and it should form part of any baseline study.

For further information contact: Conference Co-ordinator Gugulethu Charlie, SAIMM, Tel: +27 11 834-1273/7 • Fax: +27 11 833-8156 or +27 11 838-5923 E-mail: gugu@saimm.co.za • Website: http://www.saimm.co.za

The Southern African Institute of Mining and Metallurgy is proud to host the

FURNACE TAPPING 2018 CONFERENCE 15–16 October 2018—Conference 17 October 2018—Technical Tours

Nombolo Mdhluli Conference Centre Kruger National Park South Africa

BACKGROUND A controlled tapping process is essential for all pyrometallurgical smelting furnaces. Draining of metal and slag in a controlled fashion is as important as maintaining a closed tap-hole when required. From a furnace containment perspective, the tap-hole area is one of the high wear, and therefore high risk. Tap-hole failures can have devastating impacts on smelters. Risk mitigation steps therefore addresses safety and health of operators, consequential property damages, and loss of production (business interruptions). Tap-hole life-cycle design and management require development of high quality materials, equipment, and methods that not only take criteria for normal operating conditions into account, but also those of maintenance, repairs, and relines. Furnace Tapping 2014 was a first of its kind event. The event focused on the challenges associated with furnace tap-hole life-cycle design and management, and on finding ways to address these challenges, for which no miracle one-size-fits-all solution exists. South Africa, which produced 18 commodities at more than 75 sites applying smelter technology, was an ideal breeding ground for such an event. Although it had a strong local focus, people from other parts of the world also gathered to share creative solutions to problems many had in common. The SAIMM takes pride in announcing a follow-up conference, Furnace Tapping 2018, which will again be hosted in South Africa, in October 2018. The high standard of technical papers compiled in the peer-reviewed proceedings of Furnace Tapping 2014 will be maintained. The SAIMM envisage for Furnace Tapping 2018 further documentation of tapping practices by existing operators, more reviews of current operations, and descriptive case studies in which technologies available for tap-hole design, monitoring, closure, and maintenance were applied. Of special interests are feedback on research conducted in the field, and the impact of safety, health and environmental regulations from various parts of the world, on tap-hole life-cycle design and management. Looking forward to meeting you at Furnace Tapping 2018!

KEYNOTE Prof. P.C. Pistorius (Carnegie Mellon University) J.J. Sutherland (Transalloys) Dr. M.W. Kennedy (Proval Partners) Prof. J.P. Beukes (North-West University)

WHO SHOULD ATTEND The conference is aimed at delegates from the pyrometallurgical industry operating smelters or those who support them, and includes: First line, middle, and senior management Plant/production engineers Process/development engineers Refractory engineers Design engineers Photographer ©Joalet Steenkamp

Maintenance engineers Safety practitioners It doesn’t interest me who you know or how you came to be here. I want to know if you will stand in the centre of the fire with me and not shrink back. (from the poem, The Invitation, by Oriah Mountain Dreamer)

Researchers in the field

OBJECTIVES To provide an international forum for transfer of new knowledge on the design, maintenance, and operating practices surrounding the tapping of pyrometallurgical smelters, and to discuss methods and results of research conducted in the field.

FOR FURTHER INFORMATION CONTACT: Gugulethu Charlie, Conferencing Co-ordinator Saimm, P O Box 61127, Marshalltown 2107, Tel: +27 (0) 11 834-1273/7 E-mail: gugu@saimm.co.za, Website: http://www.saimm.co.za

NATIONAL & INTERNATIONAL ACTIVITIES 2017

18–20 October 2017 — 7th International Platinum Conference Platinum—A Changing Industry Protea Hotel Ranch Resort, Polokwane Contact: Gugulethu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 25 October 2017 — 14th Annual Student Colloquium Johannesburg Contact: Gugulethu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 6–7 November 2017 — Coal Preparation Society of India 5 International Conference & Exhibition ‘Coal Washing: a sustainable approach towards a greener environment’ Silver Oak Hall, India Habitat Centre, Lodhi Road, New Delhi-110003, www.cpsi.org.in

2018 25–28 February 2018 — Infacon XV: International Ferro-Alloys Congress Century City Conference Centre and Hotel, Cape Town, South Africa Contact: Gugulethu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 12–14 March 2018 — Society of Mining Professors 6th Regional Conference 2018 Overcoming challenges in the Mining Industry through sustainable mining practices Birchwood Hotel and Conference Centre, Johannesburg, South Africa Contact: Camielah Jardine

Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 11–14 June 2018 — SAIMM: Diamonds — Source to Use 2018 Conference Birchwood Conference Centre (Jet Park) Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za 8–11 July 2018 — Copper Cobalt Africa in association with The 9th Southern African Base Metals Conference Avani Victoria Falls Resort, Livingstone, Zambia Contact: Gugulethu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 13–14 September 2018 — ASM Conference 2018 ‘A focus on Illegal, traditional, customary and Small Scale Mining in Sub-Saharan Africa’ Birchwood Conference Centre, Boksburg Contact: Gugulethu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 15–17 October 2018 — Furnace Tapping 2018 Conference Nombolo Mdhluli Conference Centre, Kruger National Park, South Africa Contact: Gugulethu Charlie Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: gugu@saimm.co.za Website: http://www.saimm.co.za 2018 — Digitalization and the Internet of (Mining) Things Conference Birchwood Conference Centre (Jet Park) Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za

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17–20 October 2017 — AMI Precious Metals 2017 The Precious Metals Development Network (PMDN) Protea Hotel Ranch Resort, Polokwane Contact: Camielah Jardine Tel: +27 11 834-1273/7 Fax: +27 11 838-5923/833-8156 E-mail: camielah@saimm.co.za Website: http://www.saimm.co.za

Company Affiliates The following organizations have been admitted to the Institute as Company Affiliates

3M South Africa (Pty) Limited

Expectra 2004 (Pty) Ltd

Ncamiso Trading (Pty) Ltd

AECOM SA (Pty) Ltd

Exxaro Coal (Pty) Ltd

New Concept Mining (Pty) Ltd

AEL Mining Services Limited

Exxaro Resources Limited

Northam Platinum Ltd - Zondereinde

Air Liquide (Pty) Ltd

Filtaquip (Pty) Ltd

PANalytical (Pty) Ltd

AMEC Foster Wheeler

FLSmidth Minerals (Pty) Ltd

AMIRA International Africa (Pty) Ltd

Fluor Daniel SA (Pty) Ltd

Paterson & Cooke Consulting Engineers (Pty) Ltd

ANDRITZ Delkor(Pty) Ltd

Franki Africa (Pty) Ltd-JHB

Anglo Operations Proprietary Limited

Fraser Alexander (Pty) Ltd

Polysius A Division Of Thyssenkrupp Industrial Sol

Anglogold Ashanti Ltd

Geobrugg Southern Africa (Pty) Ltd

Precious Metals Refiners

Arcus Gibb (Pty) Ltd

Glencore

Rand Refinery Limited

Atlas Copco (SA) (Pty) Ltd

Hall Core Drilling (Pty) Ltd

Redpath Mining (South Africa) (Pty) Ltd

Aurecon South Africa (Pty) Ltd

Hatch (Pty) Ltd

Rocbolt Technologies

Aveng Engineering

Herrenknecht AG

Rosond (Pty) Ltd

Aveng Mining Shafts and Underground

HPE Hydro Power Equipment (Pty) Ltd

Royal Bafokeng Platinum

Axis House (Pty) Ltd

Impala Platinum Limited

Roytec Global (Pty) Ltd

Bafokeng Rasimone Platinum Mine

IMS Engineering (Pty) Ltd

Runge Pincock Minarco Limited

Barloworld Equipment -Mining

Ivanhoe Mines SA

Rustenburg Platinum Mines Limited

BASF Holdings SA (Pty) Ltd

Joy Global Inc. (Africa)

Salene Mining (Pty) Ltd

BCL Limited

Kudumane Manganese Resources

Becker Mining (Pty) Ltd

Leco Africa (Pty) Limited

Sandvik Mining and Construction Delmas (Pty) Ltd

BedRock Mining Support (Pty) Ltd

Longyear South Africa (Pty) Ltd

Bell Equipment Limited

Lonmin Plc

BHP Billiton Energy Coal SA Ltd

Lull Storm Trading (Pty) Ltd

Blue Cube Systems (Pty) Ltd

Magnetech (Pty) Ltd

Bluhm Burton Engineering (Pty) Ltd

Magotteaux (Pty) Ltd

Bouygues Travaux Publics

MBE Minerals SA (Pty) Ltd

CDM Group

MCC Contracts (Pty) Ltd

CGG Services SA

MD Mineral Technologies SA (Pty) Ltd

Chamber of Mines

MDM Technical Africa (Pty) Ltd

Concor Opencast Mining

Metalock Engineering RSA (Pty)Ltd

SMS group Technical Services South Africa (Pty) Ltd

Concor Technicrete

Metorex Limited

Sound Mining Solution (Pty) Ltd

Council for Geoscience Library

Metso Minerals (South Africa) (Pty) Ltd

SRK Consulting SA (Pty) Ltd

CRONIMET Mining Processing SA( Pty) Ltd

Minerals Operations Executive (Pty) Ltd

Technology Innovation Agency

CSIR Natural Resources and the Environment (NRE)

MineRP Holding (Pty) Ltd

Time Mining and Processing (Pty) Ltd

Mintek

Timrite Pty Ltd

MIP Process Technologies (Pty) Ltd

Tomra (Pty) Ltd

Modular Mining Systems Africa (Pty) Ltd

Ukwazi Mining Solutions (Pty) Ltd

MSA Group (Pty) Ltd

Umgeni Water

Multotec (Pty) Ltd

Webber Wentzel

Murray and Roberts Cementation

Weir Minerals Africa

Nalco Africa (Pty) Ltd

Worley Parsons RSA (Pty) Ltd

Data Mine SA Department of Water Affairs and Forestry Digby Wells and Associates DRA Mineral Projects (Pty) Ltd Duraset Elbroc Mining Products (Pty) Ltd eThekwini Municipality

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Perkinelmer

Sandvik Mining and Construction RSA(Pty) Ltd SANIRE Sasol Mining Sebilo Resources (Pty) Ltd Sebilo Resources (Pty) Ltd SENET (Pty) Ltd Senmin International (Pty) Ltd Smec South Africa

Namakwa Sands (Pty) Ltd

VOLUME 117

Infacon XV: International Ferro-Alloys Congress 25-28 February 2018 Century City Conference Centre and Hotel, Cape Town, South Africa

INFACON (International Ferro-Alloys Congress) was founded in South Africa in 1974 by the SAIMM (Southern African Institute of Mining and Metallurgy), Mintek (then the National Institute for Metallurgy), and the Ferro Alloys Producers' Association (FAPA) when the first INFACON was held in Johannesburg. The intention of INFACON is to stimulate technical interchange on all aspects of ferro-alloy production.

Topics for discussion:

Who should attend

Topics include but are not limited to

Metallurgists

Operational updates from ferro-alloy producers

Ferro-alloy producers

Technical aspects of ferro-alloy production Status of the ferro-alloys markets FeCr, FeMn, FeNi, FeV, FeSi, SiMn, etc. Effects of electricity cost and availability Energy efficiency and recovery Pre-treatment technologies New technologies and processes Safety Environmental issues Carbon dioxide emissions and climate change Government policies affecting ferro-alloys Carbon tax Export restrictions or subsidies

Steel and stainless steel producers Smelter operations managers Plant general managers Engineers, technicians, and scientists Process engineers Engineering companies Furnace equipment and refractory suppliers Researchers / Academics Specialists in production, economics, and the environment Policy makers Investors Students

Sustainability Use of natural gas Sale of ore versus ferro-alloy production Market supply and demand Future of the ferro-alloys industry in South Africa The impact of UG2 chromite from the PGM industry Fines, tailings, and low-grade ores Volatility of ore and ferro-alloy prices Approach to a circular economy How to extract maximum value from resources Other topics of relevance to ferro-alloy production

For further information please contact: Gugulethu Charlie • Conference Co-ordinator • E-mail: gugu@saimm.co.za Website: http://infacon15.com

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