VOLUME 115
NO. 10
OCTOBER 2015
advanced metals initiative
28–30 October 2015
our future through science
advanced metals initiative ADVANCED METALS INITIATIVE The Advanced Metals Initiative (AMI) was established by the Department of Science and Technology to facilitate research, development and innovation across the advanced metals value chain.
GOAL To target significant export income and new industries for South Africa by 2020 through the country becoming a world leader in sustainable metals production and manufacturing via technological competence and optimal, sustainable local manufacturing of value-added products, while reducing environmental impact.
To lead a global revolution in advanced metals generating significant export income and new industries for South Africa while reducing environmental impact.
STRATEGY The AMI takes an integrated approach across the entire value chain from resource development to metal production and the manufacture of end-products, to achieve its goal, through: • Reducing the energy required to produce metals by 30%; • Increasing asset productivity by 30%; • Developing technologies that can enable new industries for South Africa; and, • Reducing the full environmental impact products by 50%.
TECHNOLOGY NETWORKS The AMI’s technology networks include: • the Light Metals Development Network (LMDN) for titanium and aluminium coordinated by the Council for Scientific and Industrial Research (CSIR); • the Precious Metals Development Network (PMDN) for gold and the platinum group metals (PGMs), coordinated by Mintek; • the Nuclear Materials Development Network (NMDN) for hafnium, zirconium and monazite co-ordinated by the Nuclear Energy Corporation of South Africa (Necsa). • the Ferrous Metals Development Network (FMDN) for ferrous and base metals, co-ordinated by Mintek.
life cycle of metals
LMDN
NMDN
PMDN
FMDN
Lighter, functional alloy materials for the automotive and aerospace industries.
Beneficiation for nuclear materials used in nuclear reactors.
Value-added PGM products (Autocatalysts, PGM coatings, etc.).
Beneficiation of resources and materials solutions for the transportation, energy and petrochemical industries.
- expand the country’s technical capacity; - develop the use of metals in new applications and more diverse industries; and - develop industrial localisation.
LIGHT METALS NETWORK
DEVELOPMENT
• Global demand for ultralight, ultrastrong, recyclable metals is growing as the world switches to lowemission vehicles, energy-saving devices and sustainable products. • Aluminium demand is forecast to increase by 30% in the near future, while for the emerging industrial light metal, titanium, the sky is the limit. • For its part, South Africa has a mature aluminium industry, which is among the country’s top exporters, and one of the world’s richest titanium resources on which to build a new industry. The LMDN sees South Africa becoming a world leader in light metals. The LMDN conducts scientific research activities along the entire value chain, from resource development to primary metal production, fabrication, casting, joining technologies and manufactured products.
Enabling Platform of Projects and Partnerships
The AMI promotes collaborative research between the science councils, higher education institutions, and industry. Human resource development is critical for the networks to:
The aim is to create a globally competitive integrated light metals industry, to develop superior cost-effective technologies and manufacturing systems, and to reduce energy use, greenhouse emissions and environmental impact.
CONTACT DETAILS Llanley Simpson Tel: +27 12 843 6436 Fax: +27 86 681 0242 Cell: +27 83 408 6910 Email: llanley.simpson@dst.gov.za
NUCLEAR MATERIALS DEVELOPMENT NETWORK • The global upsurge in energy demand has led to a renewed focus on nuclear energy and related nuclear materials like zirconium and hafnium. • Zirconium is used as cladding in nuclear reactors and zirconium carbide has applications in future nuclear reactors. • South Africa has a vast resource of zircon and supplies 30% of the world market. The NMDN seeks to beneficiate zirconium and hafnium across the value chain through the preparation and purification of intermediate metal salts, metal manufacturing and optimum zirconiumalloys. The NMDN targets alternative, novel, economic and environmentally friendly manufacturing processes for the metal pair zirconium/hafnium via existing plasma and fluorochemical expertise.
advanced metals initiative FERROUS METALS DEVELOPMENT NETWORK The FMDN presents a unique opportunity to simultaneously add value to several minerals that South Africa possess in large quantities such as iron, chromium, manganese, vanadium, etc. while addressing key material challenges experienced by strategic sectors of the economy such as the transportation, energy and petrochemical industries. The FMDN R&D programmes are done within a tripartite collaborative framework involving industry, academia and science councils. The broad objectives of the FMDN can thus •be summarised as follows: Beneficiation of South Africa’s ferrous resources to stages 3 and 4 undefined.
PRECIOUS METALS DEVELOPMENT NETWORK Precious metals are characterised by their high density and cost, which make them less attractive for use in bulk components and more viable in coating/deposition technologies (chemical and physical) for various applications in which the unique properties of these high-value metals are beneficial. The PMDN assists South Africa in retaining the precious metals value matrix through the identification, research and promotion of new technologies and applications to support the long-term development of the mining industry.
The key focus is on: • creating new industries; • supporting existing industries; • localisation of technology.
supported by
• Improvement of the country’s to our capability future through science
produce high-end ferrous products, especially those that are needed by other critical sectors of the economy, such as petrochemical, energy generation, transportation, etc. • Generation of local know-how (innovation). • Human Capital Development which will alleviate the shortage in scientific and technological qualifications and skills in these sectors and thereby ensuring the sustainablity and the competitiveness of the local industry. This will also improve SA’s attractiveness as an investment destination. • Promotion of local and international collaboration in the field of ferrous metallurgy.
The Southern African Institute of Mining and Metallurgy OFFICE BEARERS AND COUNCIL FOR THE 2015/2016 SESSION Honorary President Mike Teke President, Chamber of Mines of South Africa Honorary Vice-Presidents Mosebenzi 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 President R.T. Jones President Elect C. Musingwini Vice-Presidents S. Ndlovu A.S. Macfarlane Immediate Past President J.L. Porter Honorary Treasurer C. Musingwini Ordinary Members on Council Z. Botha V.G. Duke I.J. Geldenhuys M.F. Handley W.C. Joughin M. Motuku D.D. Munro
G. Njowa A.G. Smith M.H. Solomon J.D. Steenkamp M.R. Tlala D. Tudor D.J. van Niekerk
Past Presidents Serving on Council N.A. Barcza R.D. Beck J.R. Dixon M. Dworzanowski F.M.G. Egerton H.E. James
G.V.R. Landman J.C. Ngoma S.J. Ramokgopa M.H. Rogers G.L. Smith W.H. van Niekerk
Branch Chairmen Botswana
L.E. Dimbungu
DRC
S. Maleba
Johannesburg
I. Ashmole
Namibia
N.M. Namate
Northern Cape
C.A. van Wyk
Pretoria
P. Bredell
Western Cape
A. Mainza
Zambia
D. Muma
Zimbabwe
S. Ndiyamba
Zululand
C.W. Mienie
Corresponding Members of Council 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
PAST PRESIDENTS *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)
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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)
Honorary Legal Advisers Van Hulsteyns Attorneys Auditors Messrs R.H. Kitching Secretaries 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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Editorial Board
Editorial Consultant D. Tudor
Typeset and Published by 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
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THE INSTITUTE, AS A BODY, IS NOT RESPONSIBLE FOR THE STATEMENTS AND OPINIONS ADVANCED IN ANY OF ITS PUBLICATIONS. CopyrightŠ 1978 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. U.S. Copyright Law applicable to users In the U.S.A. The appearance of the statement of copyright at the bottom of the first page of an article appearing in this journal indicates that the copyright holder consents to the making of copies of the article for personal or internal use. This consent is given on condition that the copier pays the stated fee for each copy of a paper beyond that permitted by Section 107 or 108 of the U.S. Copyright Law. The fee is to be paid through the Copyright Clearance Center, Inc., Operations Center, P.O. Box 765, Schenectady, New York 12301, U.S.A. This consent does not extend to other kinds of copying, such as copying for general distribution, for advertising or promotional purposes, for creating new collective works, or for resale.
VOLUME 115
NO. 10
OCTOBER 2015
Contents Journal Comment by J.T. Nel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
iv–v
President’s Corner by R.T. Jones . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
vii
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? by W.E. Stumpf . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
893
Friction processing as an alternative joining technology for the nuclear industry by D.G. Hattingh, L. von Wielligh, W. Thomas and M.N. James . . . . . . . . . . . . . . . . . . . . . . . . .
903
Neutron- and X-ray radiography/ tomography: non-destructive analytical tools for the characterization of nuclear materials by F.C. de Beer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
913
Non-destructive characterization of materials and components with neutron and X-ray diffraction methods by A.M. Venter. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
925
Fluorine: a key enabling element in the nuclear fuel cycle by P.L. Crouse . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
931
Titanium and zirconium metal powder spheroidization by thermal plasma processes by H. Bissett, I.J. van der Walt, J.L. Havenga and J.T. Nel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
937
Plasma technology for the manufacturing of nuclear materials at Necsa by I.J. van der Walt, J.T. Nel and J.L. Havenga . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
943
Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma operating at low to atmospheric pressures by J.H. van Laar, I.J. van der Walt, H. Bissett, G.J. Puts and P.L. Crouse. . . . . . . . . . . . . . . . . . .
949
A redetermination of the structure of tetraethylammonium meroxidotrichlorido(thenoyltrifluoroacetylacetonato- 2-O,O')niobate(V) by R. Koen, A. Roodt and H. Visser . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
957
A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form by C.J. Postma, H.F. Niemand and P.L. Crouse . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
961
Glow discharge optical emission spectroscopy: a general overview with regard to nuclear materials by S.J. LĂśtter, W. Purcell and J.T. Nel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
967
The influence of niobium content on austenite grain growth in microalloyed steels by K.A. Annan, C.W. Siyasiya and W.E. Stumpf. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
973
The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel by H.J. Uananisa, C.W. Siyasiya, W.E. Stumpf and M.J. Papo . . . . . . . . . . . . . . . . . . . . . . . . . . .
981
International Advisory Board 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 115
NO. 10
OCTOBER 2015
advanced metals initiative
All papers featured in this edition were presented at the
Nuclear Materials Development Network Conference 28–30 October 2015
our future through science
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R.D. Beck J. Beukes P. den Hoed M. Dworzanowski B. Genc M.F. Handley R.T. Jones W.C. Joughin J.A. Luckmann C. Musingwini J.H. Potgieter R.E. Robinson T.R. Stacey
Journal Comment This special edition of the Journal of the Southern African Institute of Mining and Metallurgy is dedicated to the Nuclear Materials Development Network (NMDN) of the Advanced Metals Initiative (AMI) of South Africa’s Department of Science and Technology (DST). The AMI consists of four networks: the Light Metals Development Network (LMDN) which is coordinated by the CSIR, the Precious Metals Development Network (PMDN) and the Ferrous Metals Development Network, both coordinated by Mintek, and the NMDN which is coordinated by the South African Nuclear Energy Corporation SOC Ltd (Necsa). The NMDN focuses on the development and improvement of nuclear materials in order to enhance the safety of nuclear reactors, the characterization of nuclear materials by various analytical techniques including nuclear techniques, and the beneficiation of South African minerals across the whole value chain with the aim to manufacture nuclear components in South Africa in the future. The Integrated Resource Plan for Electricity for South Africa (generally referred to as the IRP2010) identified nuclear power as an important contributor to the total electricity mix of the future. It is planned to add 9.6 GW of electricity generated by nuclear power to the national grid by 2030. This implies that six to eight new nuclear power plants (NPPs) need to be built in South Africa. Nuclear technology localization and local content will be very important factors in the realization of the nuclear new-build programme. Zirconium alloys are used as cladding material of the nuclear fuel in NPPs. All the zirconium metal that is used in this application is extracted from the mineral zircon, of which South Africa is the second-largest producer in the world. The NMDN has developed a unique plasma and fluoride beneficiation process to manufacture nuclear-grade zirconium metal powder from locally mined zircon. Hafnium is always associated with zirconium in nature, but it has to be separated from zirconium due to its high thermal neutron absorption cross-section. This property of hafnium is, however, being exploited by the nuclear industry in applications where absorbance of neutrons is required. Apart from its nuclear applications, hafnium is also used in electronics, optics, hightemperature-resistant ceramic materials, and in the aerospace industry. The NMDN has consequently also embarked on the development of hafnium products for nuclear and non-nuclear applications. The unfortunate Fukushima incident is still fresh in the minds of everybody. High temperatures and rapid oxidation of the zirconium cladding material in contact with the coolant water led to the generation of hydrogen gas and the spectacular hydrogen explosions that were witnessed by everyone on television. Since then, there has been a major drive in the nuclear industry to make zirconium alloys more resistant to oxidation at high temperatures. The application of zirconium carbide and silicon carbide layers on zirconium nuclear fuel tubes is but one of the research programmes that the NMDN is pursuing in this regard.
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Thorium is envisaged to be an important nuclear fuel for the future Generation IV high-temperature gas-cooled nuclear reactors or molten salt nuclear reactors. The mineral monazite contains a significant concentration of thorium along with valuable rare earth elements (REEs) such as neodymium (Nd), cerium (Ce), praseodymium (Pr) and yttrium (Y). Nd, for example, is a crucial ingredient in permanent magnets that are used in wind turbines to generate electricity. Monazite is a by-product from the heavy mineral sand industry in South Africa, and is also found in hard-rock orebodies, for example at Steenkampskraal and Zandkopsdrift in the Western Cape Province. Monazite, a rare earth phosphate, is extremely chemically inert. Conventional chemical extraction procedures are very harsh, environmentally unfriendly, and produce radioactive waste. The NMDN is investigating new plasma fluoride beneficiation methods to recover the REEs and to separate the contained thorium and uranium values. Human capital development forms a fundamental pillar of the AMI. In 2015, the NMDN has 28 postgraduate students enrolled at various universities in South Africa to contribute to the building of a sound local knowledge base in nuclear technology. The participation of these universities in the NMDN projects is highly appreciated. The postgraduate students of the NMDN express their extreme gratitude towards the DST for creating this programme as a launching platform for their careers. The papers that are published in this special journal of the Journal are a selection from the papers that will be presented at the 2015 academic peer-reviewed process of the AMI, of which this conference is the grand finale for postgraduate students. This year it is being hosted by the NMDN and will take place at the Nelson Mandela Metropolitan University in Port Elizabeth from 28–30 October 2015. It is noteworthy that this is the first time in the history of the AMI that the special edition of the Journal that is dedicated to the annual AMI conference will be published ahead of the conference. In this regard the Chairman thanks the organizing committee of the conference, the referees, and the SAIMM publications team for their hard work towards achieving this historic goal. The NMDN and Necsa express their sincere gratitude towards the Honourable Minister Naledi Pandor and the Honourable Director General of the Department of Science and Technology, Dr Phil Mjwara, for their continuing support of the AMI and the NMDN over many years. I hope that everyone will find the conference an enjoyable event and that this special edition of the SAIMM Journal will constitute a major contribution to the nuclear science and engineering fraternity of South Africa.
J.T. Nel NMDN Coordinator and Conference Chairman
Mining in the global village
nt’s
ide Pres
er Corn
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t has been just over fifty years since the concept of the ‘global village’ was introduced by Marshall McLuhan in 1964, yet it remains relevant to our everyday experiences. Modern communication technologies have seemingly shrunk the world even further since then. Ray Tomlinson, in 1971, was the first person to send mail from one computer to another over a network (and also initiated the practice of using the @ sign to direct the networked electronic mail message to a particular user at a particular computer). In 1997, e-mail volume overtook postal mail volume, as more and more people recognized the convenience of this almost immediate, yet still asynchronous, mode of communication. That same year, 1997, saw the registration of the google.com domain name, and searching for information was transformed for ever, as people came to rely on the Google search engine to navigate the World Wide Web (invented by Tim Berners-Lee in 1989). It seems hard to believe that it’s been only about 20 years since the mass popularization of the World Wide Web (arguably one of the world’s greatest inventions since the wheel). Nowadays, we can almost instantly read about (or watch) events happening anywhere in the world. The interconnectedness of today’s world has led to a direct link between the slowing down of the rate of growth in urbanization in China and the state of the economy in Rustenburg, for example. There is also much mobility of people between countries and continents. Many engineers trained in South Africa work in Australia, and many Australian engineers work in the USA, and so on. The SAIMM maintains strong links with similar societies in other countries. In November 2011, an inaugural meeting was held in London between several leading international mining and metallurgical societies – AusIMM (Australasian Institute of Mining and Metallurgy), CIM (Canadian Institute of Mining, Metallurgy and Petroleum), IOM3 (Institute of Materials, Minerals and Mining), SAIMM (Southern African Institute of Mining and Metallurgy), and SME (Society for Mining, Metallurgy and Exploration). The meeting was intended to foster cooperation between the various organizations, to discuss opportunities for improving and sharing benefits to members, and to benchmark the institutions against each other. Further meetings between these societies were held in September 2012 in Las Vegas (SME), in February 2013 in Denver (SME), in February 2014 in Cape Town (SAIMM), in October 2014 in Vancouver (CIM), and in March 2015 in Hong Kong (AusIMM). Agreements have been signed between these societies, resulting in the formation of what is known as the Global Mineral Professionals Alliance (GMPA). Discussions were held about the state of the mining industry in the various countries, as well as the structure and strategies of the societies represented. There was broad agreement that the societies would offer services to each other's members at member rates. This is a significant benefit to SAIMM members, as they can attend international conferences held by AusIMM, CIM, IOM3, and SME at the same cost as members of those societies. Calendars of events are circulated between the organizations to coordinate major events and minimize clashes. The flagship project of the GMPA is OneMine.org, a database of over 100 000 technical papers that is freely available to be used by the members of GMPA societies. Support of this project – both financially and by sharing technical papers – is a necessary precondition for a society to belong to the GMPA. Participating societies also agree to publicize their GMPA affiliation on their websites, and to share meeting calendars and information about each other’s international events. Representatives of each society meet once a year to exchange information, to maintain a common set of standards for technical events, and to look for further ways to increase member benefits with reciprocal arrangements. This also provides an opportunity to share approaches and resources to deal with global problems shared by all. Until asteroid mining becomes accepted practice, we will have to settle for this global approach on Planet Earth.
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R.T. Jones President, SAIMM
MINTEK – a global leader in minerals and metallurgical research and development
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a1
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? by W.E. Stumpf*
South Africa recently announced a resurgence in its commercial nuclear power programme. The implications for the development of the necessary high-level manpower within South Africa’s tertiary educational system and its national research and development (R&D) capacity in materials science and engineering, as well as in other engineering disciplines, are placed into perspective. An organized national process of developing this manpower by moving away from the previously high-risk and costly ’large programmes’ to rather a selection of ‘small and better’ research projects and a redefinition of what constitutes ‘nuclear materials’ are proposed as parts of this strategy. Keywords materials science, pressurised water reactor, zirconium based cladding materials, uranium enrichment, steam generator.
Introduction Since 2008, and more particularly in 2014/2015, South Africa has woken up to the fact that significant steps need to be taken to ensure sufficient electricity generating capacity for the future, even beyond the coal-fired stations at Medupi and Kusile currently under construction. It is, therefore, encouraging to see some active large-scale wind farms in the Eastern Cape near Jeffrey’s Bay, in the Couga area, and others in the Western Cape already in operation. In addition, many solar energy projects are also progressing from the small localized scale to larger programmes in the Northern Cape, which may contribute some capacity on a national basis. South Africa needs to tap into its renewable resources of wind and solar much more, but will these projects solve the country’s long-term industrial needs? Unfortunately not. One cannot run mines and trains on solar cells. Industry needs reliable baseload capacity, and with very limited easily accessible hydrocapacity this leaves really only coal, possibly natural gas, and nuclear power as options. South Africa’s current over-reliance of about 90% on coal–fired power, however, places it in an internationally vulnerable position, and diversification into a more equitable energy mix should be a national priority for the medium to long term. South Africa cannot simply ignore the mounting evidence of The Journal of The Southern African Institute of Mining and Metallurgy
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Broad classification of nuclear materials in a pressurized water reactor Although a modern nuclear power reactor such as a pressurized water reactor (PWR) consists
* Professor in Physical Metallurgy, University of Pretoria. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. OCTOBER 2015
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Synopsis
significant climate change confronting the human race, as the IPCC cautioned in 2013, and will have to adjust its future energy reliance to a more balanced combination of sources. After many years of international conferences, meetings, and working group sessions, the world is no nearer to finding an equitable and binding international agreement on measures to curb climate change. It is, therefore, highly unlikely that the more acceptable low-emission scenarios such as the RCP2.6 (Figure 1) are realistic, and current trends appear to indicate that the world is facing a more pessimistic climate change future, such as the RCP8.5 scenario. Does this mean that South Africa will need to completely phase out coal-fired power in the medium to long term? No, that would be impossible, and even irresponsible, but it does mean that a future energy mix of about 50% coal-fired, 25% nuclear-based, and 10% imported gas-fired power, with the remaining 15% consisting of renewable energy sources, would be a typical future to plan for. Such a scenario would constitute a baseload capacity of about 80–85% with the remainder comprising renewable energy sources, mainly wind and solar. Such a turnaround from a very high to a more reasonable dependence on coal plus a still limited nuclear dependence will place heavy demands on South Africa’s technical expertise to select, evaluate, and later to supply the materials that are ‘fit for purpose’ in the planned nuclear power programme.
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? in essence of the same main components as those for a coalfired plant, i.e. a heat source, a steam generating system, and a steam-driven turbine/generator combination, the operating and safety requirements make a typical PWR a far more complex system that requires specialized materials. Figure 2 shows a broad overview of the typical materials currently in use in a modern PWR. Note the wide range, from low-alloy steel to more sophisticated ferrous and stainless steel alloys, from nickel-based creep-resistant alloys to corrosionresistant titanium condenser tubes, from zirconium-based
fuel cladding to boron-based control rod materials, from electrically conductive copper to cathodically protected tube sheet and, last but not least, oxide fuel pellets. Production and manufacturing processes for these materials range from cast components to wrought and welded tubes and sheet, from passivated surfaces to corrosion-resistant weldcladding, from sophisticated to more conventional heat treatments, from high purity to standard material purities, from solid to porous sintered items, and so on. Such a wide range of materials of construction poses a tremendous challenge to South Africa’s materials engineers and scientists if they wish to grow into and actively participate in an expanding nuclear power programme. To simply sit back while all of the know-how is imported, even in the long term, is not an option. On the other hand, to consider actively mastering the know-how for all of the above materials is also unrealistic. Some hard choices, therefore, need to be taken to rather focus on those areas where maximum benefit can be gained within the limited research resources at the country’s disposal.
The need for materials science in PWR technology
Figure 1 – Estimated IPCC global surface temperature changes for various models of climate control through curbing CO2 emissions (IPCC, 2013)
In assessing the broad research focus areas of South Africa’s science, engineering, and technology (SET) sector in preparation for a future resurgence of nuclear power, one needs to firstly recognize the somewhat onerous process of development, testing, evaluation, and safety assessment before adoption, as described so elegantly by Hoeffelner (2011) (Figure 3). The entire cycle of materials development, from conceptual definition until final introduction in practice, can in essence be separated into two main focus areas: firstly, the
Figure 2 – Typical structural materials in use in a modern PWR (Zinkle and Was, 2013)
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Where should the national R&D in materials science fit into South Africa’s future nuclear power programme?
upper issues of technology, and secondly, the lower issues of design and safety assessment. The two focus areas go handin-hand, and South Africa’s endeavours in nuclear technology over the past three or four decades have taught some hard lessons of the consequences of focusing primarily only on the development of the technology, without planning for the resources to bring the technology into safe, reliable, and cost-effective commercial fruition, which placed the entire process at risk of termination. This was a classic technology push instead of a technology pull approach. The demands of the entire development cycle as depicted in Figure 3 can partly be recognized in the unfortunate terminations of the uranium enrichment programmes (both the Vortex and the Molecular Laser Isotope Separation (MLIS) systems) and the development of high-temperature gascooled reactor technology in the pebble bed modular reactor (PBMR) programme. In all of these programmes the technology developed by South Africa’s scientists and engineers was on an equal level with international norms, leading even to international participation in the MLIS program. During a visit in the early 1990s to the Vortex uranium enrichment Z-plant for low-level enrichment, a group of international engineers from a leading country in the area of enrichment just shook their heads and commented: ‘we would never have been able to design and build such a plant’. Why, then, did all three of these programmes falter in the end? The exact reasons are, of course, different in each case, but overall all three really faltered due to one common factor: the lack of sustainability of resources needed to take each one through the entire life-cycle of development shown in Figure 3, and then into full commercial viability. It was as simple as that! South Africa must recognize that it is a relatively small country with very limited resources, and furthermore, new technology always carries a high risk of failure even if the R&D is on a par with international norms. The crucial question regarding South Africa’s future nuclear power programme, is therefore a fundamental one: The Journal of The Southern African Institute of Mining and Metallurgy
Research focus area: technology of nuclear materials Zirconium-based cladding materials A very visible illustration of the resources required for a ‘large and new’ programme is provided by the development of improved zirconium-based cladding materials for PWR technology. VOLUME 115
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Figure 3 – The process of new materials development, testing and evaluation, design, and safety assessment in nuclear materials before acceptance and introduction (Hoeffelner, 2011)
‘Should South Africa’s materials scientists and engineers attempt to be technology leaders as in the past, or should we rather aim to be technology followers, but in the process improve on an incremental basis what others have already done, i.e. a ‘small and better’ focus? The answer to this basic question can be sought inter alia in the path taken by Japan after the end of the Second World War, when a shattered country with limited own resources had to ’climb out of its ashes’ by emulating what others had done, but doing it incrementally better. Within a decade or two, Japan had become internationally renowned for its highquality cameras, binoculars, television sets, and many other electronic and engineering goods. There is a lesson to be learnt here. South Africa should avoid ‘large and new’ high-risk technology programmes and focus rather on the ‘small and better’ technical areas that will incrementally draw South African R&D, together with local industry, into growing participation in the future nuclear power programme. A second strong argument for ‘small and better’ lies in the inherent safety and performance guarantees that have to be provided by the reactor vendor. Local participation in the supply of key components, particularly those associated directly with the so-called ‘nuclear island’, will most likely be very limited for many years to come, at least until South Africa’s industrial base has reached a level of sophistication equal to those of the reactor vendor countries. Does this, therefore, mean that no local R&D resources should be focused on these materials? No, not at all! For purposes of design evaluation, operational optimization, and safety evaluation as depicted in the lower half of Figure 3, decision-makers need to have a clear understanding of the limits of structural materials and a feeling for the behaviour of these materials under severe operational conditions, which often requires much more than ‘literature knowledge’. This route will be called ‘understanding better’. In designing a roadmap for South Africa’s materials R&D capacity in the future nuclear power programme, one can therefore once more return to the model of Hoeffelner in Figure 3 and identify two main areas that need to be addressed: ➤ Research aimed at the incremental development or improvement of the upstream technology of materials aimed at future supply into a growing nuclear power generating capacity, i.e. the ‘small and better’ route ➤ Developing an understanding of the positive and negative limits of those materials for design, life assessment, and safety evaluation purposes, i.e. the ‘understanding better’ route. Each of these will be explored in some detail, with specific examples from the nuclear industry.
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme?
Figure 4 – The development of potential advanced zirconium-based cladding materials (CEA-INSTN, 2008)
Figure 6 – In-reactor oxide thickness measurements of Zircaloy-4 (red and yellow data points) and the new Zr-Nb alloy M5 (bottom black and green data points) (CEA-INSTN, 2008)
Figure 5 – Technical challenges to be addressed in the development of any advanced zirconium-based cladding materials (CEA-INSTN, 2008)
Firstly, where would one have to focus in selecting a new development area? The leading nuclear countries of the world are committing very substantial human and research resources to the development of new alloys. The difficulty in choosing a new alloy towards which South Africa could make a meaningful contribution is immediately apparent. The development of any new alloy involves a wide range of difficult technological challenges, which often require a compromise in final properties. There is therefore an inherent risk in any choice of R&D on advanced cladding materials. Finally, the new alloy needs to be proven by means of numerous costly and lengthy in-reactor tests to evaluate its safety performance.
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Considering all of the above, it is therefore quite clear that for South Africa to embark on such a ‘large and new’ programme of cladding alloy development would be illogical, particularly since a country such as France has seemingly more than 100 engineers and scientists working on such a programme alone. Does this mean that South Africa should withdraw completely from any zirconium-based research? The answer is, clearly, ‘no!’ South Africa, together with Australia, supplies most of the world’s zirconium-based minerals, and herein lies a particular opportunity in the ‘small and better’ route. The current three major processes followed by companies in the USA, India, and France all start off with zircon (ZrO2.SiO2), which always contains small amounts of hafnium (Hf) substituted for Zr, and use various refining processes to arrive at ZrCl4 followed by reduction to hafnium-free zirconium metal in the magnesium-based Kroll process. All The Journal of The Southern African Institute of Mining and Metallurgy
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? three processes are batch operations, and all of them have environmental and safety implications. South Africa has the zirconium-based mineral resources, and is recognized internationally for its pyrometallurgical process technology. Here is a prime example of a ‘small and better’ strategy to ‘re-invent’ the whole, or only the upstream steps, of the current zirconium production routes with better technology. The incremental improvement in the three somewhat similar production routes for nuclear-grade zirconium certainly falls within the capabilities of South African scientists and engineers, and can benefit both the country and the world’s nuclear industry in the long term. Necsa, with its internationally competitive capability in fluorine and high-temperature plasma technology, is uniquely placed to accept such a challenge.
Uranium enrichment In any discussion of South Africa’s future nuclear programme, the question of uranium enrichment will inevitably arise. Should South Africa once more embark on such a ‘large and new’ venture for its future nuclear programme? This question is probably somewhat easier to answer today than some decades ago. This is due to the following reasons. ➤ South Africa’s uranium resources are found primarily in gold-bearing ores. When South Africa was one of the world’s major gold producers, it simply made sense to also beneficiate the uranium to so-called ‘yellowcake’ as a by-product. South Africa has since lost its place in the ranking of the top few gold producers. Furthermore, the fall in the price of yellowcake has made it quite uneconomical to extract the uranium from the gold recovery processes. South Africa currently (2013) accounts for only 0.9% of world uranium production (World Nuclear Association,
2013), and is superseded by African countries such as Niger (7.6%), Namibia (7.3%), and even Malawi (1.9%). Because of this unfortunate co-existence of gold and uranium, South Africa had only 5.5% of the world’s ‘Reasonably Assured Reserves’ (RAR, i.e. recoverable at a cost of less than US$130 per kilogram U), in 2009 (TradeTech, 2010), a decline from 6.5% of historical production up to 2008 ➤ Countries to the north of South Africa, however, do produce uranium as a primary product and these countries contain some noteworthy reserves. Could South Africa serve as a regional uranium enrichment centre for Africa? A select working group that was tasked by the Director-General of the IAEA in Vienna in 2007, and of which the author was a member, defined the boundaries of such regional nuclear fuel centres with one of the main criteria being ‘majority multinational control’ from outside the region, most likely with the participation of one or more of the five permanent members of the UN Security Council. To bring such a possibility into fruition, however, is fraught with a number of difficult questions, both politically and technically: • Politically, nuclear non-proliferation and uranium enrichment will always be very sensitive topics. This is not made easier by concerns about the real intentions of Iran and North Korea (both signatories of the Non-Proliferation Treaty or NPT) and those still outside the NPT, notably India, Pakistan, and Israel. With memories of the Cold War in the previous century still fresh in people’s minds and the current instabilities in many parts of the world, it is to be expected that the five permanent members of the UN Security Council, the so-called ‘haves of nuclear weapons’ within the NPT, would strongly resist any measures to spread uranium enrichment technology.
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Figure 7 – Process flow sheets for the production of nuclear-grade zirconium of (left) Wah Chang in the USA, (middle) NFC of India, and (right) Cesuz of France (CEA-INSTN, 2008)
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? • On a regional basis there is, of course the declaration of Africa as a Nuclear Weapon-free Zone, the so-called ‘Pelindaba Treaty’ originally established in 1996 and finally ratified by the 28th Member State of the African Union on 15 July 2009. This could be offered as a guarantee of the peaceful intentions of such a regional uranium enrichment centre situated in South Africa. Internationally, however, it is to be expected that serious questions will be raised as to whether such a regional treaty is ‘watertight’ against proliferation. • The next stumbling block arises from the question: ‘What happens to the depleted UF6 from interAfrican ‘imported’ uranium after enrichment?’ Technically, natural uranium imported into South Africa from any other country would be a tradeable commodity, but not the depleted UF6, which now is most likely labelled as ‘hazardous waste’. As a signatory to the Bamako Convention to ‘Control the Ban of Imports into Africa and the Control of Transboundary Movement of Hazardous Wastes within Africa’, which was signed on 30 January 1991 in Bamako, Mali and came into force on 10 March 1999, South Africa must abide by its undertaking not to allow the movement of any hazardous waste materials across international borders within Africa. South Africa, in its role of such a regional nuclear fuel centre, would therefore have to remain the host of all of the depleted uranium after the feed material had been converted to UF6 and then enriched to low enriched uranium (LEU). ➤ Finally, should all of the above political questions be somehow resolved, the technical question of ‘international cooperation’ versus ‘go-it-alone’ in selecting the technology for enrichment, needs to be understood. Here the lessons from the past should once again be
recognized. Yes, South Africa’s SET capacity could, in principle, once more go down that road with centrifuge technology, but it will always be a high-risk strategy with the very real possibility of another failure on commercial grounds. The low-risk strategy of using the currently most competitive uranium enrichment centrifuge technology of URENCO under a licence agreement, as France has done for some years, needs to be considered. Considering all of the above, it seems that a possible South African re-entry into the area of uranium enrichment would be faced with almost insurmountable hurdles, and would need to be very carefully analysed before it is even considered.
The ‘small and better’ approach for South African SET in materials research Recognize development trends in commercial power reactor trends The development of commercial power reactor technology has progressed a long way towards the ‘Generation IV’ (GEN IV) light water reactor (LWR), with enhanced economy and safety, minimal waste generation, and, last but not least, increased proof against proliferation. South Africa is one of the participating countries in GEN IV and is, therefore, well placed to use this association in planning its nuclear power reactor programme for the future. Note the key targets of better economy, enhanced safety, less waste, and proliferation resistance. In each of these areas, South Africa’s SET capacity can certainly make a ‘smaller and better’ contribution.
Should South Africa focus its nuclear materials research on in-core or ex-core components? In considering this question for LWR technology, one inevitably focuses on UO2 and zirconium-based cladding
Figure 8 – -The Generation IV LWR development path (US Department of Energy, 2002)
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Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? Some typical incremental advances with a high impact in the LWR nuclear industry
material, but any modifications, even those that are only modest, without ultimate proof of performance under irradiation conditions, would be almost a futile exercise (Figure 6). Is the commercial transfer of in-core technology from a reactor vendor into South Africa an option? Here again, lessons from the past need to be recognized, as in the transfer in the 1980s of Koeberg fuel manufacturing technology from France in Necsa’s former BEVA programme. Although the transfer was technically successful (the BEVA fuel elements in Koeberg performed on par with those in France,) the transfer of technology incurred high costs over a period of about four years. At that point, however, the reactor vendor had already moved on with its next, more advanced fuel element design, which would have required another significant investment for South Africa to ‘stay in the game’, a situation that would recur every few years. The message here is ‘don’t even consider investing heavily in the local manufacture of critical items in the nuclear island unless a significant percentage of South Africa’s electricity generating capacity will be nuclear-based, thus warranting such an investment’. This raises the obvious question of whether South Africa really needs a reactor such as SAFARI I at all. The answer to this question is an overwhelming YES. South Africa, as one of the top three medical isotope producers in the world, should not relinquish its position. This was achieved under difficult conditions and the replacement of the ageing SAFARI I reactor needs to take that into account. Should ‘SAFARI II’ then be only an isotope-producing reactor? This would be a very unfortunate retrograde step, as the growing use of SAFARI I for non-nuclear industrial tests such as residual stresses and texture formation in metals, as well as neutron research, constitutes a powerful training and research instrument for South Africa’s national SET institutions. In the ‘small and better’ strategy, South Africa’s SET capacity should, therefore, rather focus on ex-core components with the general aim of assisting local industry to participate in a meaningful manner in the future nuclear construction programme, but always focusing on the overriding aims of the Generation IV LWR.
Watanabe (1974) once had an ‘eureka’ moment when he changed the relationship defined by the well-known HallPetch equation, that a reduction in grain size leads to the ‘holy grail’ of higher strength with improved ductility, by asking what would happen if we were to change not the size, but the nature of the grain boundaries. This opened up many studies towards understanding so-called coincident site lattice (CSL) boundaries and how to increase their percentage in a mixture of low- and high-angle grain boundaries (LAGBs and HAGBs), twin boundaries (TBs), and then CSL boundaries. CSL boundaries are high-angle boundaries but possess the special characteristics of LAGBs. The percentage of CSL boundaries can be measured with little difficulty by many modern scanning electron microscopes fitted with an electron backscatter diffraction (EBSD) capability. This innovation is now being applied to the vexing problem of general intergranular stress corrosion cracking (IGSCC) in the tubing of steam generators of PWR stations.
Figure 9 – The allowable operating stresses for a design life of at least 300 000 hours for three steels typically used in LWR technology. SA 508 Grade 3 is a low-carbon manganese-molybdenum steel ’optimized for the nuclear industry. X20 (2.25Cr-Mo steel) and P91 (X10CrMoVNb9-1) are both used for high-temperature steam piping (Buckthorpe, 2002)
Figure 10 – 100 000-hour creep rupture strength and temperature of 9–12%Cr steels in general over time with the introduction of P91 steel by the late 1980s (Klueh and Harries, 2001)
The steels used for high-temperature steam piping for pressure vessels in both conventional coal- and nuclearpowered generating stations still require better understanding in terms of properties such as creep behaviour, corrosion, weldability, and end-of-life assessment. For instance, the weldability of P91 (nominally a 9%Cr-1%Mo-V-Nb steel) in the Medupi power station required detailed attention to meet the design requirements, and for application in a nuclear power station the welding codes will even be stricter. The steady improvements in the service performance of P91 (Figure 10) shown below is evidence of ’small and better’ improvements over time, achieved only through dedicated research.
Innovative materials science in solving a stress corrosion cracking (SCC) problem through grain boundary engineering (GBE)
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The development of advanced steels for pressure vessels and steam piping
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme?
Figure 11 – A schematic cutaway showing the internals of a PWR steam generator (Staehle, 2007) and general corrosion problems experienced in LWR systems with the IGSCC Alloy 600 used in the U-tubes in the PWR steam generators (top of the figure) since the inception of commercial PWR technology (Palumbo, 1993)
Figure 12 – Alloy 600 (Ni –15.74Cr – 9.1Fe) sensitized for 1 hour at 600°C followed by 120 hours of corrosion testing according to ASTM G28 with (i) its conventional microstructure and (ii) after grain boundary engineering (Lin et al., 1995)
In a number of groundbreaking patents and publications (Palumbo, 1993; Aust, Erb, and Palumbo,1994; Palumbo, Lehockey, and Lin, 1998; Was, Thaveeprungsriporn, and Crawford, 1998), Palumbo and others have found the means to increase the density of CSL boundaries in Alloy 600 (a Ni16Cr-9Fe alloy) through iterative strain-annealing or iterative strain-recrystallization, typically from less than 10% CSL boundaries in the unprocessed alloy to as high as almost 50% in the iteratively strain-annealed or strain-recrystallized form. Note the very clear micrographic evidence of less grain boundary penetration from the surface for a grain-boundaryengineered Alloy 600 that has been subjected to laboratorysimulated stress corrosion cracking. Note that in the coarse-grained microstructure in Figure 14, the increase in the CSL boundary density from 20% to 34% was sufficient to dramatically lower the alloy’s creep rate at 360°C, while in a fine-grained Alloy 600 the improvement in creep strength appears to be even greater. In both cases,
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the improvements in creep strength were achieved with a relatively modest increase in CSL boundary density, with no further improvements at higher CSL densities.
Shifting the focus from ‘technology’ to ’technology plus design and safety’ Up till now, the focus was very much on the technology of nuclear materials, but as Hoeffelner (2011) has shown (Figure 3), this is barely half the story necessary for a resurgence of a national nuclear power programme. The other half will require a significant body of high-level SET capacity with better understanding to technically evaluate offers from vendors, ask the right questions, retrieve the required design and performance data, and critically compare differences in design approach from the point of view of: ➤ Expected performance ➤ Life assessment The Journal of The Southern African Institute of Mining and Metallurgy
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? ➤ The demonstrated safety of the system being offered under all possible external or internal scenarios before an operating licensing by the NNR (National Nuclear Regulator) can even be considered. In the short term this can be achieved by hiring in nuclear consultants (preferably with no national affiliation with the vendorcountry), but this cannot be sustained in the long term for reasons of cost. This approach would also represent a lost opportunity for developing that high-level manpower required for firstly the pre-operational evaluation phase of each offer, then the construction and operational stage of a number of power stations, and finally waste management and the eventual decommissioning of the power stations. Building up such a body of high-level manpower may appear to be a daunting task, but it is no accident that Hoeffelner (Figure 3) places the technology focus at the point of entry for design and safety assessment. For instance, a
Figure 13 – Dependence of total grain boundary cracked fraction on CSL boundary fraction in Alloy 600 (Ni-16Cr-9Fe) (with varying amounts of carbon in solution and for the case of grain boundary carbides) for strains of 15% and 20% with testing in 360°C PWR primary circuit water at a strain rate of 3 ×10-7 per second (Alexandreanu, Capell, B., and Was, 2001)
postgraduate engineer or scientist who has gained experience in one aspect of the arc welding of zirconium-based cladding material for a Masters or a PhD degree will have absorbed the intricate nature of this alloy far more than in a few months of reading in a library or attending postgraduate courses, without going through the demanding process of an in-depth literature review, research planning and execution, reporting of results, and finally modelling these results in an informative discussion; with all of this being critically reviewed at the end by one or more external examiners from the nuclear industry. In addition, international publications that are peer-reviewed by experts in the field add to demonstrating ‘better understanding’. However, a basic change will, have to occur in the definition of which materials are nuclear and which are not, as Figure 2 has shown. Nuclear materials in a resurgent nuclear power programme are not simply limited to zirconium, uranium, and a very few others while the rest are seen as ‘conventional’ and therefore technically outside the current narrow definition of ‘nuclear materials’. A steam generator on a PWR can never be viewed as simply another type of heat exchanger, and thereby not warranting the level of regulatory supervision that a nuclear reactor pressure vessel does (and even the latter material is traditionally not viewed, in South Africa at least, as a ‘nuclear material’). For reasons of public safety and acceptance/assurance, the process of obtaining an ASME Section III N-certificate on equipment used in a nuclear reactor are far more onerous than if the same equipment is used in a conventional power station. This places most of the materials listed in Figure 2 in an entirely new class, that should be dealt with appropriately. Some typical areas in the category of ‘better understanding’ of an advanced PWR to consider in research could include: ➤ Projects that entail the welding of various components, such as fuel cladding, pressure vessel steel, steam piping, steam generator items, etc. ➤ Projects that entail the strength/ductility/creep/fracture relationships of the above materials at room temperature as well as at typical steam operating conditions ➤ Corrosion properties, including IGSS, pitting corrosion,
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Figure 14 – Constant load creep data on Alloy 600 (Ni-16Cr-9Fe) of a (i) coarse-grained and (ii) a fine-grained microstructure in solution annealed and grain boundary engineered conditions. The creep tests were carried out under argon at 360°C with creep stresses of 300 and 450 MPa respectively. In (iii), the dependence of the steady-state creep rate and the percentage of cracked boundaries on the fraction of CSL boundaries on the coarse-grained material is shown (Was, Thaveeprungsriporn, and Crawford, 1998)
Where should the national R&D in materials science fit into South Africa’s future nuclear power programme? hydriding of zirconium-based cladding, iodine SCC of zirconium alloys, SCC of steam piping and steam generator components, etc.
IPCC. 2013. Climate Change 2013. The Physical Science Basis. IPCC Working Group 1 Contribution to the Fifth Assessment Report of the Intergovernmental Panel on Climatic Change. Stocker, T.F., Qin, D., Plattner, G.-K., Tignor, M., Allen, S.K., Boschung, J., Nauels, A., Xia, Y., Bex, V., and Midgley, P.M. (eds).AR5, WGI 5th Assessment report.
Conclusions The following aspects need to be incorporated into a national manpower training programme that should be well under way long before the first reactor vendor’s technical offer is received. ➤ Set up a Governmental Advisory Board consisting of the main role-players in an expanded nuclear power programme. These would include Eskom, the NNR, Necsa, senior representatives from industry, public representatives, and the like. The constitution of such a body would, however, have to be very carefully drafted to totally exclude items not associated purely with training of high-level manpower, as any other issue raised in such a forum could compromise the independence of the NNR in ruling on safety issues ➤ Establish initially at least one (and later possibly a second or third) research centre at a South African university that will undertake the postgraduate training and development necessary for scientists, engineers, and technologists required by the programme. This needs to be done in close association with Necsa ➤ Ensure that the local R&D effort eschews end-use nuclear materials, but rather focuses on upstream processes in a ‘small and better’ focus ➤ Expand Necsa’s mandate to include research and development on materials covered by the broader definition of nuclear material. The Advanced Metals Initiative of DST can play a deciding role in overseeing healthy cooperation between the Ferrous Metals Development Network managed on behalf of the AMI by Mintek and the Nuclear Metals Development Network managed on behalf of the AMI by Necsa.
KLUEH, R.L., and HARRIES, D.R. 2001. High chromium ferritic and martensitic steels for nuclear applications. Monograph 3. ASTM International, West Conshohocken, PA.
LIN, P., PALUMBO, G,. ERB, U., and AUST, K.T. 1995. Influence of grain boundary character distribution on sensitization and intergranular corrosion of Alloy 600. Scripta Metallurgica et Materialia, , vol. 33. p. 1387.
PALUMBO G. 1993. Thermomechanical processing of metallic materials. Int. Patent Application PCT/CA1993/000556.
PALUMBO, G., LEHOCKEY, E.M., and LIN, P. 1998. Journal of Metals, February. p. 40.
PWR STEAM GENERATORS https://www.google.co.za/search?biw=1366&bih=673&noj=1&site=webh p&source=hp&q=pwr+steam+generators&oq=PWR+Steam+Generators&gs _l=hp.1.0.0i30.8889.22694.0.25409.20.19.0.1.1.0.642.4064.2-13j51.14.0.msedr...0...1c.1.62.hp..5.15.4080.OZJadVVd4AE [Accessed 6 March 2015].
STAEHLE, R.W. 2007. Anatomy of proactivity. International Symposium on Research for Aging Management of LWR and its Future Trends. 15th anniversary of the Institute of Nuclear Safety Systems Inc., Fukui City, Japan, 22–23 October 2007.
TRADETECH. Not dated. World Uranium Production and Requirements.
Acknowledgments The views expressed in this presentation are the author’s own and do not necessarily reflect official policy of the University of Pretoria or of the South African Government. The author would also like thank many colleagues at Necsa and the University of Pretoria for valuable input and comments. This contribution is published with the permission of the University of Pretoria.
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ALEXANDREANU, B., CAPELL B., and WAS, G.S. 2001. Combined effect of special grain boundaries and grain boundary carbides on IGSCC of Ni-16Cr-9FexC alloys. Materials Science and Engineering, vol. A300, p94–104. AUST, K.T., ERB, U. and PALUMBO, G. 1994. Interface control for resistance to intergranular cracking. Materials Science and Engineering.A, vol. A176, p.329. BUCKTHORPE, D. 2002. https://odin.jrc.ec.europa.eu/htr-tn/HTR-Eurocourse2002/Buckthorpe_582.pdf [Accessed 6 March 2015] CEA-INSTN. 2008. Zr in the Nuclear Industry. CEA-INSTN Seminar, April 2008. HOEFFELNER, W. 2011. Materials databases and knowledge management for advanced nuclear technologies. Journal of Pressure Vessel Technology, vol. 133, no. 1. pp. 014505 1-4. doi:10.1115/1.4002262.
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misorientation effects on creep and cracking in Ni-based alloys. Journal of Metals, February, pp. 44–49.
WATANABE, T. 1974. An approach to grain boundary design for strong ductile polycrystals. Res Mechanica, vol.11. p. 47.
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a2
Friction processing as an alternative joining technology for the nuclear industry by D.G. Hattingh*‡, L. von Wielligh*, W. Thomas† and M.N. James‡
The process of joining materials by friction is based on generating the heat necessary to create a solid-state mechanical bond between two faying surfaces to be joined. In simple terms, the components to be joined are subjected to frictional heating between rubbing surfaces, causing an increase in interface temperature and leading to localized softening of interface material, creating what is described as a ’third body’ plasticized layer. This plasticized zone reduces the energy input rate from frictional heating and hence prevents macroscopic melting. The plasticized layer can no longer transmit sufficient stress as it effectively behaves as a lubricant (Boldyrev and Voinov, 1980; Godet, 1984; Singer, 1998; Suery, Blandin, and Dendievel, 1994). The potential for this solid-state frictional joining process to create high-performance joints between, for example, dissimilar materials with limited detrimental metallurgical impact, and reduced defect population and residual stress level, has had a very significant impact on fabrication and repair in industrial sectors such as transport. This paper presents a brief overview of the advances made within the family of friction processing technologies that could potentially be exploited in the nuclear industry as alternative joining and repair techniques to fusion welding. Modern friction processing technologies can be placed into two main categories: those that make use of a consumable tool to achieve the intended repair or joint (friction stud and friction hydro-pillar processing) and those making use of a non-consumable tool (friction stir welding). The most mature friction joining technology is friction rotary welding, where a joint is formed between original parent materials only. A new addition in this category is linear friction welding, which opens the potential for joining complex near-net-shape geometries by friction heating. The continuous innovation in friction processing over the last 25 years has led to the development of a number of unique processes and applications, highlighting the adaptability of friction processes for specialized applications for high-value engineering components. Keywords friction processing, joining, leak sealing, weld repair technology.
Introduction Assessing the potential for using alternative joining and repair techniques for specific applications in a given industrial sector is informed by an appreciation of the current size and potential growth in that sector. Power generation by nuclear fission has been out of favour over the past few years, due to several high-profile nuclear incidents, e.g. the Fukushima Daiichi nuclear disaster, as well as concerns over disposal of nuclear waste and spent fuel rods. However, due to rapid increases in the demand for electricity, coupled The Journal of The Southern African Institute of Mining and Metallurgy
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with many fossil-fuel fired plants reaching end-of-life, there is renewed interest in installing additional nuclear capacity. Proposing friction processing as an alternative (and relatively untried) joining technology for the nuclear industry might be viewed as potentially perilous because of stringent requirements for validation of weld and repair procedures. The intention in this paper is to introduce some modern developments in the friction processing arena and to outline the potential that these processes hold for manufacturing of new components and for maintenance and life extension of ageing nuclear power plant. The World Nuclear Association (WNA, 2015) states that as at April 2015, there were 437 nuclear reactors in operation, with another 65 under construction and a further 481 either in planning or proposed. Clearly, there is significant potential for the application of friction processing techniques in this power sector. The age of reactors varies, from firstgeneration reactors developed in the 1950s to the current generation III reactors introduced in 1996 in Japan, with generation IV reactors still under development with a proposed introduction date of 2020. The main reactor types currently operational are pressurized water reactors (PWRs) – 273 units, boilingwater reactors (BWRs) – 81 units, pressurized heavy-water reactors (CANDU-PHWR) – 48 units, gas-cooled reactors (AGR and Magnox) – 15 units, light-water graphite reactors (RBMK and EGP) –15 units, and fast neutron breeder reactors (FBR) – 2 units (WNA, 2015). The majority of reactors currently in operation are relatively elderly in terms of their design life and would require decommissioning
Friction processing as an alternative joining technology for the nuclear industry in the short term if life extension methodologies are not adopted. The high capital cost and extensive lead time involved in bringing new nuclear plants online make the whole idea of accurately determining the remaining life of existing plants very topical. The graphs in Figure 1 and Figure 2 were populated from data published by the WNA (2015) in 2015 and give an indication of the current operational plants, plants under construction, and proposed plants. Considering Figure 1, which shows the percentage contribution from nuclear fission, over the decade 2003 to 2013, to the electric power needs of both the ‘BRICS’ nations and of developed countries, it is clear that most of these countries made very little capital investment in additional nuclear generation capacity, with most countries showing a moderate decline in nuclear contribution. With the exception of France, and to a lesser extent Brazil, most countries currently obtain less than 10% of their electricity needs from nuclear. The future for nuclear power looks much more promising when one considers the potential additional capacity from the current build programme, and the planned and proposed additions to the nuclear fleet over the next 20 years (Figure 2). Therefore, taking an overview of the installed, planned, and proposed future nuclear capacity, and factoring in developments in respect of the safer, more standardized and cost-effective generation IV reactors, there is considerable opportunity to plan for the implementation of alternative joining and repair technologies in the nuclear industry. The basic principles of generating electricity from nuclear fission are very similar to those used in fossil-fired thermal power plants, if viewed from the secondary circuit (steam
generation) point onwards. Steam generated by controlled nuclear fission is used to drive turbines, which are interconnected with electrical generators. A number of the main materials used in the construction of nuclear power plants, including Zircaloy, can be joined or repaired by friction-based processes. The joining of thin-wall sections intended for use in fuel rod application by rotary friction processes is currently being studied at the Nelson Mandela Metropolitan University (NMMU) Work is also being done on Cr-Mo-V alloy steels and various grades of stainless steel.
Friction processing technologies One of the simplest ways to warm up your hands is to rub them together. This action generates heat in proportion to the rapidity of the rubbing motion, the pressure applied on the rubbing surfaces, and the duration. Similar concepts are used to provide the primary thermal energy source in friction processing (FP), allowing lower temperature solid-state joints and repair processes to be applied to a wide variety of materials (particularly the magnesium and aluminium light alloys) used in manufacturing industry (Thomas and Duncan, 2010). In summary, friction processing comprises a set of solid-state joining techniques that are carried out by generating heat at a contact interface. This softens the material and allows plastic flow to occur which, in combination with an applied pressure, can be used to create strong, low-defect, and low tensile residual stress bonds (Maalekian , 2007; Gibson, 1997; ASM International, 1993; Serva-Tech Systems, n.d. (a); Thomas and Dolby, 2001). There are various types of friction processes, as illustrated in Figure 3. The main friction processing joining technologies considered in this paper are rotary friction welding (RFW), friction stir welding (FSW), friction hydro-pillar processing (FHPP, which can employ either tapered or parallel sided configurations), and linear friction welding (LFW). LFW is a
Figure 1 – Nuclear power share for the period 2003 to 2013 (WNA, 2015)
Figure 2 – Nuclear power generation landscape expressed in MWe (WNA, 2015)
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Figure 3 – Friction process technologies (TWI) (Thomas and Dolby, 2002) The Journal of The Southern African Institute of Mining and Metallurgy
Friction processing as an alternative joining technology for the nuclear industry relatively new and versatile process that permits the joining of irregular shapes, allowing for near-net-shape fabrication. In this range of friction technologies, there are autogenous processes that form a joint by using parent material only (RFW and LFW), i.e. without the addition of any filler metal, processes that make use of a consumable tool (FHPP), and processes that use non-consumable tools (FSW). There are four basic classifications for relative motion in friction welding (BSI Group, 2000; Serva-Tech Systems, n.d. (b)), namely: rotary, angular oscillation, linear reciprocation, and orbital motion. Figure 4 illustrates these various motions. Fabrication by frictional heating is a well-established technology. The first use of friction knowledge to process and shape materials dates back to 1891, when Bevington (1891) used heat generated by friction to form and join tubes.
Rotary friction welding (RFW) RFW is a well-established ‘workhorse’ friction joining technology that has been widely used for over a century in numerous applications ranging across the automotive, construction, aerospace, and medical sectors. Two alternative RFW techniques, known as continuous drive and inertia friction welding, are currently in use; the latter is the most widely adopted process. In both techniques joints are made by placing a rotating component in contact with a stationary component while applying a load perpendicularly to the contact interface. Once the interface reaches the appropriate welding temperature to plastically displace and fuse the materials by forging, rotation is stopped and the weld is allowed to solidify while maintaining the forging force (Thomas and Duncan, 2010; Thomas, Nicholas, and Kallee, 2001) as illustrated in Figure 5. The differences between the two methods are mainly in process control, with continuous
drive welding being done with a constant rotational speed that may be varied at different stages of the weld cycle, while during inertia welding the process begins at a relatively high rotation speed, which gradually reduces to zero (Thomas and Duncan, 2010).
Friction hydro-pillar processing (FHPP) Like most of the modern friction technologies, FHPP (Figure 6) was invented and patented by TWI (World Centre for Materials Joining Technology) in the UK in the early 1990s. A description of the process, as given by Thomas and Nicholas (1992), states that friction hydro-pillar bonding is achieved by rotating a consumable tool coaxially in a hole while under load. The frictional heat generated results in the production of a pillar of continuous plasticized layers until the hole is filled. The plasticized material consists of a series of shear layers or interfaces which solidifies under the influence of the applied force. One of the main original observations was that the plasticized material advances more rapidly than the axial feed rate of the consumable tool, which results in the rising of the frictional interface along the consumable tool to form the dynamically recrystallized deposit material. During the process a good degree of plasticization must be maintained at the interface, allowing hydrostatic forces to be transmitted axially and radially to the inside of the hole in order to achieve a good metallurgical bond (Thomas and Nicholas, 1992). A large body of reported work from various laboratories has shown that good mechanical integrity can be achieved in FHPP welds, with metallographic examination indicating a fine-grained microstructure in the welded zone (TWI Bulletin, 1997). The original geometry proposed for the FHPP process was based on a horizontal cylindrical arrangement, with specific clearance between the tool and the sides of the hole. However, in more recent applications tapered holes and consumables are becoming more common, in which the angle of the two tapers is slightly different (Thomas and TempleSmith, 1996; Bulbring et al., 2012; Meyer, 2001). Wedderburn (2013) reported that the tapered arrangement allows for an additional reactive force to develop horizontally from the sidewall in addition to the vertical hydrodynamic force, which can be utilized for heat generation when making the joint, and that this is beneficial for materials with poor extrusion properties.
Figure 4 – Relative motion classification in friction welding (Wedderburn, 2013)
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Figure 5 – Schematic of the basic rotary friction welding process (Pinheiro, 2008)
Friction processing as an alternative joining technology for the nuclear industry A number of research institutes have claimed that good quality FHPP welds can be made in steel and certain nonferrous materials using a parallel hole geometry. In this context a ‘good quality’ weld is characterized by high impact, tensile, and bend properties (TWI Bulletin, 1997). This sentiment is supported by Wedderburn (2013) Bullbring, et al., (2012) who reported that good mechanical properties were achieved with a taper configuration on most thick-wall steam pipe materials evaluated. Where welds were made using parallel tool and holes, periodic changes in the microstructure were observed, where regions within the deposit material were not fully transformed. This effect was more noticeable when a high tool rotation speed and consumable displacement (upset) rate were used. Microstructural variations were also exacerbated when the hole depth to hole diameter exceeded a ratio of about 3.5:1 (Meyer, 2003). There is limited published work on the influence of process parameters on FHPP welds. From the work by Meyer (2003), Wedderburn (2013), and Bulbring et al., (2012) on the use of FHPP for high-strength low-alloy steel, it is apparent that hole geometry is more critical than tool shape. Meyer (2003) also noted that the influence on heat generation and bond quality is similar to that observed in conventional friction welding, although the rotational speed required for a good FHPP weld is significantly lower than in conventional friction welding. The forging force, which has a major influence on weld properties in conventional friction welding, appears to have little or no influence on most FHPP welds and its effect is limited to the upper area near the surface region (Wedderburn, 2013; Bullbring, et al., 2012; Meyer, 2003).
Friction stir welding (FSW) FSW must be considered as one of the leading innovations in solid-state joining in the last 50 years. Invented and patented by TWI (1991), FSW is a joining technique that allows both low and high melting point metals, including aluminium, lead, magnesium, steel, titanium, zirconium, and copper, to be continuously welded with a non-consumable tool (Nicholas, 1998; Reynolds, Seidel, and Simonsen, 1999; Dawes, 2000). The wear rate and cost of refractory tooling is a limiting factor in welding steel alloys. One of the key benefits of this solid-state friction welding technique is the
Figure 7 – Principle of friction stir welding (Thomas, Nicholas, and Kallee, 2001)
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capability of welding dissimilar metals as well as joining certain metals that are difficult to weld by fusion processes, e.g. 2xxx series aluminium alloys. The process can best be described as a continuous ‘hot-shear’ process involving a non-consumable tool that is responsible for mixing and forging the metals across the joint line (Delany et al., 2005). Selection of the tool geometry and material are important considerations in achieving a sound joint. The most important aspects of tool design are the pin and shoulder geometry, as these are responsible for forming the joint. The basic principle of the FSW process involves plunging a rotating tool between abutting faces of the work piece and then traversing it along the joint line (Figure 7). Rotary motion of the tool generates frictional heat at the contact surfaces, softening the material and creating a plasticised zone around the tool pin and below the shoulder. Apart from generating heat, the shoulder also contains the plasticized material in the joint zone, which provides a forging action that assists with the formation of a defect-free solid-state joint. In the nuclear industry, FSW has been applied to the manufacture of copper canisters for encapsulating nuclear waste (Andersson and Andrews, 1999) as well as for leak sealing or material reprocessing techniques used in the repair of surface-breaking or near-surface defects (Von Wielligh, 2012).
Linear friction welding (LFW) (LFW is a joining technology that is ideal for welding nonsymmetrical components. This solid-state joining process generates frictional heat through the linear forced interaction between a stationary surface and a reciprocating surface (Nentwig, 1995; Nicholas, 2003). The process requires a large force applied perpendicular to the weld interface and, as there is no containment of the hot plasticized material, a continuous layer of flash is expelled during the rapid linear movement. Since this softened ’third-body’ material is not contained, surface oxides and other impurities are ejected together with the plasticized ‘flash’ (Bhamji et al., 2011). Unfortunately, LFW requires a major capital investment due to the complexity and size of the welding platforms. The high process force must be contained within the platform
Figure 8 – Schematic diagram of the linear friction welding process (Bhamji et al., 2011) The Journal of The Southern African Institute of Mining and Metallurgy
Friction processing as an alternative joining technology for the nuclear industry
Applications of friction technologies in the nuclear industry A number of studies have shown that friction welding can be used to make sound joints and cost-effective repairs in the nuclear industry. These include the work done in a Swedish programme that evaluated FSW for encapsulating nuclear fuel waste (Andersson and Andrews, 1999), while TWI demonstrated that thermocouple probes can be fitted to boiler header domes forged from 316 austenitic stainless steel for the Hartlepool and Heysham nuclear power stations (TWI Global, 2014). Hy-Ten, an independent rebar and accessory supplier, has gained approval for supplying friction welded rebar couplers to be used in concrete construction by the nuclear industry (Maguire, 2012). Locally, the NMMU has developed two technologies with potential applications in the nuclear industry; one relating to leak sealing by FSW that employs a low-force partial-penetration methodology, and the other the registered Weldcore® process, which employs FHPP principles to core and repair sites used for creep assessment as part of remaining life evaluation.
Nuclear waste storage encapsulation using FSW This project formed part of a Swedish programme that evaluated new ways of encapsulating nuclear fuel waste before storing it underground. The capsules were manufactured from copper and consisted of a copper cylinder with a base and lid (Andersson and Andrews, 1999). The capsule demand was estimated at 200 per year. Important considerations in fabrication included inspection of the welds and guarantees on achieving defect-free joints that would prevent water entering the canister during long-term storage underground (Swedish Nuclear Fuel and Waste Management Co. (SKB), 2001). Typical canister dimensions were diameter 1050 mm, length 4830 mm, and a total weight of 27 t. To satisfy the design requirements for environmental and chemical resistance, a copper alloy equivalent to EN 133/63 with a 50 mm wall thickness was selected for the canisters. To improve creep resistance of the copper, 50 ppm of phosphorus was specified in the alloy (Swedish Nuclear Fuel and Waste Management Co. (2001). The project developed an electron-beam welding solution for sealing the capsules and The Journal of The Southern African Institute of Mining and Metallurgy
also considered the use of FSW. The main consideration in selecting FSW as a potential alternative process was the solid-state nature of the joining process, together with the ability to make high-quality welds with a fully automated platform and standardized tool technology. The investigation of FSW as an alternative joining process started by considering the welding of 10 mm thick copper plate strips about 0.5 m in length. This study was used to find a suitable tool material as well as evaluate different tool geometries. The only reference study at the time involved FSW of thick-section aluminium plate with melting temperatures (659°C) well below that of copper (1083°C). However, copper tends to soften at lower a temperature, which allows the use of FSW on thick copper sections (Swedish Nuclear Fuel and Waste Management Co., 2001). This study led to the specification of process forces and power requirements for welding 50 mm copper plate. This was done by progressively increasing the plate thickness to 50 mm while checking that the weld microstructure remained acceptable for the identified application. The weld zone showed a fine equiaxed grain structure in the weld nugget region. Based on the success of these initial trials, SKB proceeded with the development of a FSW platform to weld lids to short (2 m) sections of 50 mm thick tubes. The tube was placed in a vertical position with the welding head and tool in a horizontal position while the welding head rotated around the tube. For each weld, specific weld parameters were recorded, providing a complete record of all the essential variables and their control during the welding process. These results, in conjunction with non-destructive testing, form an important resource for interpreting weld parameters and data. Once adequate circumferential friction stir-welded joints were achieved, attention was focused on eliminating the exit hole left at the end of the weld run. This project has clearly demonstrated the feasibility of fabricating capsules by FSW to enclose nuclear waste for long-term storage.
Leak sealing by FSW Nuclear facilities situated close to the ocean could experience material degradation, especially of stainless steel vessels and pipework, arising from environmentally-induced stress corrosion cracking (SCC). This prompted an investigation into identifying a suitable refurbishment technique that would provide a cost-effective and permanent leak-sealing/SCC repair technique. Implementation of the repair process was required to have no influence on the operation of the plant, and also to fulfil the stringent nuclear safety regulations. The NMMU was therefore requested to investigate the feasibility of using FSW as a potential leak-sealing technique for 304L stainless steel. This study focused on establishing a technology procedure for the identified application, which entailed reviewing the process, tool material, and tool geometry designs to develop a solution suitable for hightemperature FSW of unsupported, partial penetration leak sealing of 304L stainless steel with water backing. The approach taken in the investigation was initially to establish a theoretical processing parameter window from published literature. Subsequently, a suitable tool alloy and tool profile were identified before preliminary experimental work led to a low-force processing window for 304L stainless steel. Various experiments were carried out to determine the effect of process parameters on weld quality and the VOLUME 115
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structure, while achieving the required level of displacement and process control necessary to ensure joint integrity and geometrical accuracy means that tight specifications are set on all system components and control algorithms. Hence LFW is considered feasible for the production of high-value engineering components, e.g. joining of aero-engine compressor blades to compressor disks (Wanjara and Jahazi, 2005). LFW applications are predicated on the suitability of the process for joining materials with good high-temperature properties, low thermal conductivity, and acceptable compressive yield and shear strength properties (Wanjara and Jahazi, 2005). Low thermal conductivity helps confine heat to the interface region, while appropriate hightemperature properties (high yield strength and melting point) allow a high level of frictional heat generation to occur before plastic collapse. Materials with good high-temperature mechanical properties and low thermal conductivities, e.g. titanium, zirconium, and nickel alloys, are particularly suitable for LFW applications.
Friction processing as an alternative joining technology for the nuclear industry processing forces required to achieve high quality. The governing factor was found to be the downwards forging force required to achieve a fully consolidated, partial penetration weld that would seal a leak in stainless steel sheet resulting from, for example, SCC. Finally, a complete weld procedure specification was proposed and evaluated by welding on a mock-up of the real application with water seeping through the simulated crack during the welding process. In principle, this is a similar concept to FSW under water, which has been reported by Ambroziak and Gul (2007). Such applications require refractory tool materials, and W-25wt%Re was selected for the leak-sealing application as it provides good resistance under conditions of uneven surfaces, high vibratory loads, and plate deflection (Von Wielligh, 2012). Additional benefits include the ability to reprofile the tool using standard machining processes, hence increasing tool life. Process development focused on establishing a low-force processing window for the 304L
Figure 9 – Difference in the Z-force feedback for different control strategies (Von Wielligh, 2012)
stainless steel that minimized plate deflection during welding in an unsupported condition. The investigation led to the development of a new dual-control tool plunge strategy, in which the operator controls the maximum downwards Z-force during the plunge stage while simultaneously controlling the plunge rate and depth (see Figure 9). Plunge rate is, however, also a function of the Z-force feedback, which could be seen as a significant advantage because the maximum plate deflection at the start of the weld can therefore be limited (Von Wielligh, 2012). The minimum Z-force required for repeatable weld consolidation using tools with a 14 mm, 12 mm, or 10 mm diameter shoulder was found to be 14.25 kN, 10.5 kN, and and 9 kN respectively (Von Wielligh, 2012). This investigation showed that water cooling and water seepage through the simulated SCC had only a minor effect on weld quality and that the dominant factor influencing weld quality was plate deflection. A full preliminary welding procedure specification was developed for both the 14 mm and 12 mm shoulder diameter tools, but only the procedure for the 14 mm tool was deemed robust enough for industrial application. Un-etched images of typical weld cross-sections obtained during the repair trials of SCC in stainless steel are shown in Figure 10. This figure shows the defect population observed at the start, middle, and end positions in two welds, A and B, made using different weld parameters while water was seeping through the SCC, using two different sets of process parameters. On the etched specimen in the macrograph in Figure 11, the lighter etching material clearly shows a well-defined flow pattern in the weld cross-section. This material was found to be tungsten, a by-product of tool wear that has diffused into the weld. The darker line extending beyond the flaw intended to simulate a stress corrosion crack at the bottom of weld cross-section is termed a ‘lazy S’ defect in FSW and is not part of the simulated crack. It is generally ascribed to the
Figure 10 – The defect population found in weld cross-sections along the weld length representing the start, middle, and end of welds for two different welding parameter sets (Von Wielligh, 2012)
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Friction processing as an alternative joining technology for the nuclear industry presence of oxidation and corrosion by-products of the wire cutting process, and serves as a good indication of how actual corrosion products are likely to be distributed across the weld. Vickers microhardness tests did not show any significant change in the hardness of the weld area or heat-affected zone compared with the parent material. This is to be expected, as austenitic stainless steels are not hardenable by heat treatment (Von Wielligh, 2012). The feasibility of applying this FSW process to crack repair in industry was tested using a mock-up of the industrial application of a tank containing a stress-corrosion crack. A simulated crack was successfully welded in an unsupported 304L stainless steel plate surface in the annealed condition with water backing up to a depth of 2 mm. The weld repair process required making a number of overlapping friction stir welds, with a step distance between them equal to the pin diameter. The step direction was towards the retreating side of the weld, because flash formation arising from plate deflection made it impractical to step towards the advancing side. The typical weld repair pattern is shown in Figure 13. The first weld, W1, is carried out from left to right with a clockwise spindle rotation. The retreating side of the weld is thus located at the bottom. The second weld, W2, is carried out from right to left with an anticlockwise spindle rotation such that the retreating side remains at the bottom of the weld. This process of alternating welds is continued until the cracked area has been covered.
Figure 13 shows an example of such a weld repair with five overlapping welds, covering an area of 120 mm × 20 mm. In this figure, all the loosely attached flash has been removed. This investigation has shown that leak sealing of stresscorrosion cracks in an unsupported tank surface is possible by using overlapping, partial penetration, friction stir welding. The investigation also highlighted the fact that due to material thinning during the FSW process, the total area than can be repaired has to be limited, based on allowable plate thinning. A reduction in the original plate thickness occurs with each successive overlapping weld. The decrease in plate thickness is likely to be a function of plate thickness; in other words, dependent on constraint. Furthermore, FSW tool exit holes can be partially eliminated by using tools with a retractable pin. In summary, FSW is a potential repair technique for SCC, although for unsupported welding careful consideration must be given to the influence of plate thickness, surface condition of the area to be processed and its size, and deflection during welding, as well as the permanent deformation (thinning) induced by the process. FSW is also likely to alter the metallurgical and mechanical properties in the processed area, and this aspect also needs consideration. Associated friction processes such as friction surface dressing or friction cladding could also be considered for such repair procedures if it is accepted that the main purpose of the process is not to recover structural integrity, but rather top-seal seepage of the liquid inside the tank through stress corrosion cracks. Friction cladding is an additive process, which has the potential to reduce plate deflection and permanent deformation imposed by the friction process. Figure 14 shows the heat generation in the tool while traversing along a repair zone with water leaking through the simulated crack. The installation of the developed platform is shown in Figure 15.
Figure 11 – Macrograph of a typical friction stir processed SCC, etched with Marbles reagent [30]
Figure 13 – The visual appearance of a friction stir processed tank surface area with the loosely connected flash removed on the advancing side of the weld (Von Wielligh, 2012)
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Figure 12 – The welding pattern of five consecutive welds used to friction stir process an area of 120 × 20 mm (Von Wielligh, 2012). (Welds W1, W3, and W5 were welded from left to right and welds W2 and W4 from right to left)
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Figure 15 – Set-up for the repair process feasibility investigation, showing the FSW platform attached to a curved stainless steel water-filled test coupon (Von Wielligh, 2012)
Weldcore® sample removal and repair technology Ongoing assessment of the remaining life of hightemperature and -pressure (HTP) components in nuclear power stations is of paramount importance in ensuring their safe and cost-effective operation. Although failure can have severe consequences, economic considerations are requiring operators to extend the operating lives of plants currently in operation. To ensure safe operation, more direct experimental assessments have to be made of remaining life to complement and underpin prediction models. Neutron irradiation embrittlement is one of the concerns for the nuclear industry, as this could limit the service life of reactor pressure vessels. Improved understanding of the underlying causes of embrittlement and better calibration of prediction models will provide both regulators and power plant operators with improved estimates of vessel remaining life (Odette and Lucas, 2001). The NMMU, in collaboration with Eskom, has developed the patented Weldcore® process, which is a sampling and repair process that is currently used in fossil fuel power stations to assist in sample removal for creep damage assessment. The process involves removal of a cylindrical
Figure 16 – Illustration of four Weldcore® coring stages (Hattingh et al., 2012)
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metallurgical sample and hole preparation for repair in four steps, as illustrated in Figure 16. The hole is then repaired using friction taper hydro-pillar processing (FTHPP) as illustrated in Figure 17. Current Weldcore® applications are primarily in a coring and repair procedure for local metallurgical sample removal sites. The advantages of the process are that the metallurgical core can extend much closer to the centre of a thick-walled pressure pipe, and that the coring and repair can be accomplished during plant operation. The core hole may be either tapered or vertical-sided, depending upon the thickness of the material to be welded and the depth of the hole. For deeper repair welds, a cylindrical configuration for the hole
Figure 17 – Schematic of the FTHPP process with macro appearance of a cross-section of a weld (Hattingh et al., 2012) The Journal of The Southern African Institute of Mining and Metallurgy
Friction processing as an alternative joining technology for the nuclear industry and rotating tool is recommended to ensure that the process force and torque are limited. The consumable rotating tool is typically made of the same alloy as the base metal, although it is possible to achieve a matched, over-matched, or undermatched weld zone. On completion of the weld, the remaining portion of the consumable rotating tool can be cut off and ground flush with the metal surface. Since the ’weld’ does not produce a liquid weld puddle, the orientation of the weld is not affected by gravity, hence making the process positionindependent. The main factors making the Weldcore® process attractive to the nuclear fraternity relate to platform size (which is compact and hence suitable for on-site work) and a high degree of process automation. The Weldcore® platforms were initially developed for high-temperature steam pipes used in power stations. However, the success of the platform has led to further developments for applications that include turbine disc and reducer sections. Figure 18 shows the modular arrangement and adaptability of the welding platforms that were developed at the NMMU for specific applications of the Weldcore® process in the power generation industry. The in-situ application on steam pipes required the design and development of a portable reconfigurable platform that allowed operation at various locations throughout a power station. The modular arrangement of the welding platform, involving a frame, spindle, and drive motor, facilitates ease of transport, positioning, and mounting. Electrical and hydraulic power and control connections are operated via a remote control unit. The process and portable platform allow for the in-situ removal of a representative sample of metal from a structure (pipes, turbine rotors, etc.) that can be used for accurate metallurgical assessment of the remaining life of the structure. It provides superior data to other current methods as the excised core is a far more representative sample of the through-thickness than other techniques, with a shorter
acquisition downtime. The process has been tested on steam pipes and turbine components in the power generation industry and is in the final phases of being tested on structures on a petrochemical plant. The main benefit of the Weldcore® process is a reduction in the risk of failure and hence better management of safe life in high-value engineering components.
Discussion The introduction of a new or alternative joining and repair processes into the nuclear industry entails a number of challenges from the engineering and regulatory points of view. Nuclear power, for obvious reasons, is a tightly regulated industry, but the authors propose that by introducing improved life assessment standardization, the risk and complexity of safely operating these types of facilities could be further reduced. Cost-effective condition monitoring is an implicit part of life extension for ageing plant, and of the through-life cost and performance optimization of new installations. Friction processing presents a viable alternative methodology for some aspects of condition monitoring and repair of thermal power plant, with significant advantages accruing from its solid-state nature (which opens the door to welding of dissimilar metals). A solid-state joining process limits microstructural phase transformations, while lower process temperatures result in a significant reduction in thermal stresses and associated distortion compared with conventional fusion welding processes. The lower thermal gradient results in a narrow heat-affected zone with a fine equiaxed weld nugget and limited grain growth in the thermomechanically transformed region. These microstructures are normally associated with good strength and toughness. As illustrated in the applications presented in this paper, friction processing lends itself to a high degree of automation and process control, giving repeatable welds and allowing sophisticated monitoring of weld quality. Friction
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Figure 18 – Developmental progression of Weldcore® equipment (Hattingh et al., 2012)
Friction processing as an alternative joining technology for the nuclear industry welding offers the power generation industry a clean, reconfigurable welding process suitable for complex and bespoke applications and which can greatly reduce process time and cost of condition monitoring and repair.
References AMBROZIAK, A. and GUL, B. 2007. Investigations of underwater FHPP for welding steel overlap joints. Archives of Civil and Mechanical Engineering, vol. 7, no. 2. pp. 67–76. ANDERSSON, C-G. and ANDREWS, R.E. Fabrication of containment canisters for nuclear waste by friction stir welding. Proceedings of the 1st International. Symposium on Friction Stir Welding, Rockwell Science Center, Thousand Oaks, California, 14–16 June 1999. ASM INTERNATIONAL. 1983. Metals Handbook. Volume 6 – Welding Brazing, and Soldering. 9th edn. Materials Park, Ohio. BEVINGTON, J. (1891. Spinning tubes mode of welding the ends of wire, rods, etc, and mode of making tubes. US patent 463134, 1891. BHAMJI, I. PREUSS, M., THREADGILL, P.L., and ADDISON, A.C. 2011. Solid state joining of metals by linear friction welding: a literature review. Materials Science and Technology, vol. 27, no.1. pp. 2–12. BOLDYREV, R.N. and VOINOV, V.P. 1980. Possible reasons for the formation of extremum of torque during heating in friction welding. Welding Production, no. 1. pp. 10–12.
NICHOLAS, E.D. 2003. Friction processing technologies. Welding in the World, vol. 47, November. pp. 11–12. ODETTE, G.R. and LUCAS, G.E. 2001. Embrittlement of nuclear reactor pressure vessels. Journal of Materials (JOM), vol. 53, no. 7. pp. 18–22. PINHEIRO, G.A. 2008. Local reinforcement of magnesium components by friction processing: determination of bonding mechanisms and assessment of joint properties. Dissertation, Vom Promotionausschuss der Technischen Universität Hamburg-Harburals. REYNOLDS, A.P., SEIDEL, T.U., and SIMONSEN, M. 1999. Visualization of material flow in an autogenous friction stir weld. Proceedings of the 1st International Symposium on Friction Stir Welding, Rockwell Science Center, Thousand Oaks, California, 14–16 June 1999. SERVA-TECH SYSTEMS. Not dated (a). Friction Welding Process Basics – report 1. http://www.frictionwelding.com/report1.htm SERVA-TECH SYSTEMS. Not dated (b). Direct vs Inertia Friction Welding – report 4. Not dated. http://www.frictionwelding.com/report4.htm SINGER, I.L. 1998. How third-body process affects friction and wear. MRS Bulletin, June. pp. 37–40. SUERY, M., BLANDIN, J.J., and DENDIEVEL, R. 1994. Rheological behaviour of two phase superplastic materials. Materials Science Forum, vol. 170/172. pp. 167–176.
BSI GROUP. (2000. Welding – friction welding of metallic materials. British Standard BS EN ISO 15620-2000. London, UK.
SWEDISH NUCLEAR FUEL AND WASTE MANAGEMENT CO. 2001. Development of fabrication technology for copper canisters with cast inserts. Technical Report TR-02-07. Stockholm, Sweden.
BULBRING, D.L.H., HATTINGH, D.G., BOTES, A, and ODENDAAL, D.H. 2012. Friction hydro pillar process as an alternative repair steel structures. FERROUS 2012. Ferrous and Base Metals Development Network Conference, Magaliesburg, South Africa, 15–17 October 2012. Southern African Institute of Mining and Metallurgy, Johannesburg.
THOMAS, W. and DUNCAN, A. 2010. Friction based technology for joining and material processing – an introduction. International Conference on Welding and Joining of Materials (ICWJM), Pontificia Universidad Católica del Perú, Lima, Peru, 9–11 August 2010.
DAWES, C.J. 2000. Faster and faster - welding speed increases with tool development - one of a series of steps. Bulletin, vol. 41, no. 4. pp. 51–55.
THOMAS, W.M. 1997. Gas additions boost friction performances. TWI Connect Press. The Welding Institute, Cambridge, UK. www.twi.co.uk.
DELANY, F., LUCAS, W., THOMAS, W., HOWSE, D., ABSON, D., MULLIGAN, S., and BIRD, C. 2005. International Forum on Welding Technologies in Energy Engineering. Shanghai, China, 21–23 September 2005.
THOMAS, W.M. and DOLBY, E. 2002. Friction stir welding developments. 6th International Conference on Trends in Welding Research, Pine Mountain, Georgia. USA, 15–19 April 2002.
GIBSON, S.W. 1997. Advanced Welding. Chapter 12 – Friction Welding. McMillan, London.
THOMAS, W.M. and NICHOLAS E.D. 1992. Friction hydro pillar processing TWI. Leading Edge. TWI Connect Press. www.twi.co.uk
GODET, M. 1984. The third-body approach: a mechanical view of wear. Wear, vol. 100. pp. 437–452.
THOMAS, W.M. and TEMPLE-SMITH, P. 1996. Friction plug extrusion. UK Patent Application, GB 2306365. The Welding Institute, Cambridge, UK.
HATTINGH, D.G., DOUBELL, P., SCHEEPERS, R., NEWBY, M., VON WIELLIGH, L., ODENDAAL, D., and WEDDEBURN, I.N. 2012. Novel core sampling technique for HP turbine rotor remaining life study. 10th Electric Power Research Institute Conference, Florida, USA.`
THOMAS, W.M., NICHOLAS, E.D., and KALLEE, S.W. 2001. Friction based technologies for joining and processing. TMS Friction Stir Welding and Processing Conference, Indianapolis, Indiana, USA, November 2001.
MAGUIRE, A. 2012. Hy-Ten celebrates nuclear approvals for friction welded couplers in concrete construction. http://www.sourcewire.com/news/73407/hy-ten-celebrates-nuclearapprovals-for-friction-welded-couplers-in-concrete# MEYER, A. 2003. Friction hydro pillar processing - bonding mechanisms and properties. Dissertation, Von der Gemeinsamen Fakultät für Maschinenbau und Elektrotechnik der Technischen Universität Carolo-Wilhelmina zu Braunschweig als. Dissertation angenommene Arbeit. GKSSForschungszentrum Geesthacht GmbH. MAALEKIAN, M. 2007. Friction welding - critical assessment of literature. Journal of Science and Technology of Welding and Joining, vol. 12, no. 8. pp. 708–729. NENTWIG, A.W.E. 1995. Untersuchungen zum linear-reibsscheissen von metallen. Schweissen und Schneiden, vol 47, no. 8. pp. 648–653. NICHOLAS, E.D. 1998. Developments in the friction stir welding of metals. 6th International Conference on Aluminium Alloys, ICAA-6, Toyohashi, Japan. Sato, T., Kumai, S., Kobayashi, T., and Murakami, Y. (eds). Japan Institute of Light Metals, Tokyo.
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TWI BULLETIN. 1997. The need for gas shielding - positive advantages for two friction processes. The Welding Institute, Cambridge, UK. pp. 84-88. TWI GLOBAL. 2014. Friction welding proves perfect in nuclear work. http://www.twi-global.com/news-events/case-studies/friction-weldingproves-perfect-in-nuclear-work-097/. VON WIELLIGH, L.G. 2012. Un-supported friction stir processing of stress corrosion cracks in a 304L stainless steel tank surface. Technical Report TS008-A. eNtsa, Nelson Mandela Metropolitan University. WANJARA, P. and JAHAZI, M. 2005. Linear friction welding of Ti-6Al-4V: processing, microstructure, and mechanical-property inter-relationships. Metallurgical and Materials Transctions A, vol. 36A, no. 8. pp. 2149–2164. WEDDERBURN, I.N. 2013. Development of a creep sample retrieval technique and friction weld site repair procedure. PhD thesis, Nelson Mandela Metropolitan University, Port Elizabeth, South Africa. WORLD NUCLEAR ASSOCIATION (WNA). 2015. http://www.worldnuclear.org/info/Facts-and-Figures [Accessed 2 May 2015].
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Neutron- and X-ray radiography/ tomography: non-destructive analytical tools for the characterization of nuclear materials by F.C. de Beer*†
A number of important areas in nuclear fuel cycle, at both the front end and back end, offer ideal opportunities for the application of nondestructive evaluation techniques. These techniques do not only provide opportunities for non-invasive testing of e.g. irradiated materials, but also play an important role in the development of new materials in the nuclear sector. The advantage of penetrating radiation used as probe in the investigation and testing of nuclear materials makes X-ray and neutron radiography (2D) and tomography (3D) suitable for various applications in the total nuclear fuel cycle. The unique and different interaction modes of the two radiation probes with materials provide several opportunities. Their complementary nature and non-destructive character makes them most suitable for nuclear material analyses, analytical method development, and the evaluation of the performance of existing nuclear material compositions. This article gives an overview of the X-ray and neutron radiography/tomography applications in the field of nuclear material testing, and highlights a few of the success stories. Several selected areas of application in the nuclear fuel cycle are discussed to illustrate the complementary nature of these techniques as applied to nuclear materials. Keywords neutron radiography, X-ray radiography, SAFARI-1; non-destructive testing.
Introduction During the development of new materials for the nuclear industry, materials testing and characterization are of the outmost importance to maintain safety standards and reliability. No compromise on safety in the workplace in any area within the total nuclear fuel cycle can be tolerated, and therefore most in-situ material characterization and testing is conducted by certified and qualified personnel schooled in destructive and non-destructive testing (DT and NDT) methods. Certification and qualification in NDT can be obtained through many training centres in South Africa in accordance with European, American, and other international standards (SGS, n.d.; SAIW, n.d.; African NTD Centre, n.d.). The testing of new methods and materials related to the nuclear fuel cycle is essential for the continuing development and safety of nuclear-related materials and processes. These fundamental research initiatives are mostly performed at the laboratory scale by material and instrument scientists, researchers, and most likely postgraduate students. The Journal of The Southern African Institute of Mining and Metallurgy
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* Radiation Science Department, Necsa, Pelindaba. † School of Chemical and Minerals Engineering, North-West University, Potchefstroom. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. OCTOBER 2015
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Davies (2000) describes the role and value of NDT during maintenance and in-service inspection of nuclear power plants during outages, and particularly the monitoring of material degradation to prevent failure. Ultrasonic testing (UT), magnetic testing (MT,) and electrical testing (ET) play a major role as NDT methods for monitoring materials degradation in-situ, while atomic- and nuclear physics-based methods such as positron annihilation, neutron diffraction, as well as Xray and neutron tomography are limited to laboratory-scale experimentation. However, conventional film-based X-ray and gamma-ray radiography (RT) techniques are being applied throughout many areas of material testing in the nuclear fuel cycle. The ‘nuclear fuel cycle’ refers to the entire range of activities associated with the production of electricity from nuclear fission, entailing (International Atomic Energy Agency, n.d.): ➤ Mining and milling: from mined uranium to yellowcake ➤ Conversion: from yellowcake to gas ➤ Enrichment: increases the proportion of the fissile Isotope ➤ Deconversion: depleted uranium. ➤ Fuel fabrication: UO2 pellets – fuel pins – fuel elements ➤ Electricity generation: fuel burn-up ➤ Storage: spent fuel ➤ Reprocessing: spent fuel ➤ Radioactive waste: safe storage. The nuclear fuel cycle includes the ‘front end’, i.e. preparation of the fuel, the ‘service period’ in which fuel is used in the reactor to generate electricity, and the ‘back end’, i.e. the safe management of spent fuel, including reprocessing and reuse and disposal.
Neutron- and X-ray radiography/tomography The general nuclear fuel cycle is schematically depicted in Figure 1, showing the various activities in the production of energy through the nuclear fission process. Every activity requires conventional NDT techniques to be conducted to maintain the safe working and operation of the plants and facilities. The standard NDT methods applied to e.g. inspection of welds in piping, are (Willcox and Downes, n.d): ➤ Radiography testing (RT) ➤ Magnetic particle crack detection (MT) ➤ Dye penetrant testing (PT) ➤ Ultrasonic flaw detection (UT) ➤ Eddy current and electromagnetic testing (ET). This paper does not focus on the so-called conventional NDT techniques and their application in the nuclear sector, but rather on the non-conventional NDT techniques that are used as needed, and which constitute important research tools. In particular, penetrating radiation probes as realized in radiography/tomography are described with specific applications in material research. Quantitative and/or qualitative data obtained through applying these novel techniques in a laboratory environment adds value to many areas within the nuclear sector. The following specific activities, ranging from mining the ore to security of the waste generated, and where radiography and tomography are applied, are highlighted in this paper: ➤ Mining and geosciences: quantification of ore deposits ➤ Fuel fabrication: development and testing of new materials ➤ Electricity generation: fuel rod performance ; postirradiation examination (PIE) ➤ Radioactive waste: safe storage, civil engineering.
Analytical methods based on penetrating radiation Information about the internal structures of objects, for example the hydrogen content of Zr cladding, can be obtained by destructive analytical methods, e.g. cutting a the fuel rod in a 2D plane for analysis by electron microscopy, or sieve analysis for particle size distribution of a soil sample. In most cases, once the sample has been destroyed, no other analytical tests are possible and the larger picture (volumetric hydrogen distribution and particle size distribution in the soil) is lost. More valuable, unique, and in some cases more accurate results can be obtained only when three-dimensional information is available. For research purposes the most acceptable way to obtain information while maintaining the sample integrity is to apply a non-destructive test using penetrating radiation (either with X-rays, gamma rays, or neutrons). It is worthwhile to mention that the neutron, the fission product of the nuclear fuel cycle, can be used as a probe to investigate the integrity of the nuclear fuel itself. After irradiation the physical condition of fuel pellets, while still intact in the fuel pin, can be obtained only by means of radiography. This manner of non-invasive investigation keeps the sample intact, leaves the sample in its original form, and it is possible for other tests to be conducted subsequently on the real sample as if it was not touched. The non-invasive process allows for the generation of valuable qualitative information. However, when digital data is transformed into three-dimensional tomographic data (Banhard, 2008), it is possible to obtain high-resolution quantitative information of the internal structures and properties of the object. An example is the volumetric pore size distribution of voids within nuclear-encapsulating concrete matrixes, as well as their physical distribution throughout the sample (McGlinn et al., 2010). X-ray, gamma-ray, and neutron radiation are attenuated (absorbed and scattered by the sample) according to an exponential law (De Beer, Middleton, and Hilson, 2004): [1] where I is the intensity of the transmitted radiation beam, I0 the intensity of the incident radiation beam, μ the attenuation coefficient (cm2/g) of the material under investigation for the specific radiation type, ρ the density of the sample (g/cm3), and x the thickness of the sample (cm). The attenuation coefficient μ expresses the total attenuation, due to both the scattering and capture processes for the incident radiation. The term μρ is also called the total absorption coefficient of the sample. We assume that the quantity μρ is linearly related to its constituents: [2]
Figure 1 – Schematic depiction of the nuclear fuel cycle (Pixshark, n.d.)
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where μi is the radiation attenuation coefficient of constituent i, ρi the density of constituent i, Vi the volume fraction of constituent i, and ∑ the summation symbol for the ith components. Composites of materials will thus have a different radiation attenuation property from the individual elements. The parameters I(E), I0(E), and μ(E) reflect radiation energy dependency. This dependency, in a radiography The Journal of The Southern African Institute of Mining and Metallurgy
Neutron- and X-ray radiography/tomography context, means that materials will attenuate different radiation types by different magnitudes and thus will yield different radiological images, and also that an element has different attenuation properties at different energy levels for the same radiation type. This implies that an element can be transparent to fast neutrons (MeV energies) but can be detected easily using thermal neutrons (eV energies). An example is the thermal neutron scattering and absorption of hydrogen and boron (Domanus, 1992). Although the basic interaction of X-rays and neutrons with the elements differs, the principle of conducting radiography to obtain a two-dimensional radiograph/image of the sample is the same. Figure 2 schematically illustrates the basic components and layout of a radiography facility. A source of radiation emits penetrating radiation towards a sample. For example, a sample contains either a defect, an inclusion of another material, or a void that is abnormal for the sample, or an area that differs completely in terms of composition from the basic matrix of the sample, which results in a lower or higher density at the location within the sample. The incident radiation will be attenuated (scattered and/or absorbed) differently due to the abnormality. A sensitive area detector, with a high quantum efficiency for the detection of the specific type of penetrating radiation, registers the difference in attenuated radiation that has passed through the sample. The 2D data (image) obtained by
Figure 2 – Principle and layout of a 2D radiography set-up. A similar set-up is used for tomography, with the sample rotating in the radiation beam (Domanus, 1992)
the detector is called a radiograph and contains the integrated radiation transmitted information for the total sample in a certain orientation with respect to the source and detector configuration. The information captured in the radiograph differs in principle for X-ray and neutron radiography/tomography. The two probes are mostly utilized within the nuclear fuel cycle as non-destructive techniques in research and for nuclear material qualification and quantification. These principles are discussed in more detail in the following paragraphs.
X-ray radiography X-ray interaction with materials depends on the density of the sample – i.e. the electron cloud density (Banhard, 2008). The area detector registers a two-dimensional image (radiograph) of the object representing the internal structure density. Elements with low electron densities are not easy to resolve in a radiograph, but they are easily penetrated to reveal denser materials embedded within the sample matrix. Figure 3 presents the different X-ray attenuation coefficients (cm-1) for 125 kV X-rays for the full spectrum of elements in the periodic table. It clearly shows the increase in absorption of X-rays (darker shading) at higher atomic numbers.
Neutron radiography The interaction of neutrons with materials is totally different to that of X-rays, since neutrons, being neutral particles, interact only with the nucleus of the atom. Neutrons are not affected by even a dense electron cloud, e.g. of a lead atom (μρn = 0.38 cm-1). The thermal neutron attenuation coefficients depicted in Figure 4 shows a totally different, and in some instances an opposite attenuation capability (grey scale) to that for X-rays (Figure 3). Hydrogen, as a highly attenuating material (μρn = 3.44 cm-1), will be easy to detect and clearly visible on a neutron radiograph when embedded in e.g. a ZrTM (μρn = 0.29 cm-1) fuel pin, which is nearly transparent to neutrons. A radiograph with low- to intermediate-energy X-rays is possible as the ZrTM tube (μρX = 2.47 cm-1) attenuates most of the X-ray radiation and with no photons remaining, the H (μρX = 0.02 cm-1) cannot be registered/detected on the X-ray radiograph.
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Figure 3 – Periodic table with X-ray attenuation coefficients of the elements for 125 kV X-ray energies (Grünauer, 2005)
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Figure 4 – Periodic table with thermal neutron energy attenuation coefficients of the elements (Grünauer, 2005)
Tomography
Mining
The word ‘tomography’ comes from the Greek words ‘to cut or section’ (tomos) and ‘to write’ (graphein) (Banhard, 2008). Tomography is also known as computer tomography (CT) or computer assisted tomography (CAT) as in diagnostic investigations in the medical field. For the purpose of this article, the following semantics are adopted: CT in general description, XCT for X-ray computer tomography, and NT for neutron tomography. CT is a radiographic inspection method that uses a computer to reconstruct an image of a crosssectional plane (slice) through an object (ISO 15708-1). The resulting cross-sectional image is a quantitative map of the linear radiation attenuation coefficient, μ, at each point in the plane. The linear attenuation coefficient characterizes the local instantaneous rate at which the incident radiation is attenuated during the scan, by scatter or absorption, as it propagates through the object. To obtain this ‘map’, the sample is radiographed and thus projection data is gathered from multiple directions through many angles of the sample. For the purpose of this article, no detailed description of the 3D reconstruction process of the sample is presented. To put it simply, multiple 2D-projections are fed into a dedicated computer with a specialized computer algorithm to create cross-sectional planes of the sample. When these cross-sectional planes are stacked together, a full virtual three-dimensional image (tomogram) of the sample can be viewed and analysed.
X-ray-, gamma-ray, and neutron tomography have demonstrated their potential in the earth sciences as important diagnostic tools to generate volumetric data on geological compositions, especially advances in the area borehole core investigations, as depicted in Figure 5. This aspect is being explored further with optimum resolution obtained through the application of micro-focus X-ray tomography, as CT complements conventional destructive analytical thin-sectioning of drill core samples (De Beer and Ameglio, 2011). The raw material for nuclear fuel is uranium, which is a relatively common element that can be found throughout the world. Uranium is present in most rocks and soils, in many rivers, and in seawater. Uranium is about 500 times more abundant than gold and about as common as tin. The largest producers of uranium are currently Australia, Canada, and Kazakhstan, with Namibia rated 5th and South Africa 11th globally (World Nuclear Association, 2015). The
Application of radiography and tomography within the nuclear fuel cycle In each of the following sectors of the nuclear fuel cycle, radiography and/or tomography have been applied using either X-rays or neutrons. In some areas the application was pioneered in the early second half of the 20th century by means of film techniques. The techniques applied and the description thereof is not within the scope of this article. However, the outcomes of these film-based investigations and the results obtained will be described, together with the recent digital methods in this field.
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Figure 5 – Left: the grouped pyroxene mineral (coloured red) can be clearly seen in norite (top), and different concentration of feldspar are observed in anorthosite (bottom) in a thin slice. Right: transparent corresponding neutron tomograms. The minerals shown are only pyroxenes that are present in both norite and anorthosite but at different concentrations [archives from Necsa’s Nrad and MIXRAD system] The Journal of The Southern African Institute of Mining and Metallurgy
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Figure 6 – X-ray CT-tomogram slice showing the distribution of uraninite grains in the quartz-pyrite matrix and the carbonaceous matter (Sebola, 2014) The Journal of The Southern African Institute of Mining and Metallurgy
technique cannot be used as an independent tool for mineral characterization, but rather in support of existing mineralogical techniques. However, valuable information is added into the nuclear value chain via 3D-CT. No sample preparation is needed other than cutting a small piece of rock for analysis. The sample integrity is maintained due to the non-destructive nature of the technique, as it provides 3D information for the total volume of the sample, including internal components. Results can be obtained within about 1 hour scan at high resolution with up to 2000 projections, providing resolution down to 6 μm (including reconstruction).
Enrichment Enriched uranium is uranium in which the concentration of U-235 has been increased through the process of isotope separation. U-235 is the only nuclide existing in nature (in any appreciable amount) that is fissile with thermal neutrons (OECD Nuclear Energy Agency, 2003). Natural uranium is 99.284% U-238 isotope, with the U-235 isotope constituting only about 0.711%. The very first application of neutron radiography was in the early 1960s for nuclear fuel characterization using the film technique. The use of neutron radiography in the monitoring of isotopic enrichment in fuel pellets loaded into a fuel pin has been demonstrated by Frajtag (n.d.) (Figure 8). Gosh, Panakkal, and Roy (1983) investigated the possibility of monitoring plutonium enrichment in mixedoxide fuel pellets inside fuel pins using neutron radiography as early as 1983. Recently, Tremsin et al., (2013) investigated the very large difference in the absorption crosssections of U-235 and U-238 isotopes, as shown in Figure 9, and deduced that very accurate non-destructive spatial mapping of the enrichment level of fuel pellets, as well as of the distribution of other isotopes in the spent fuel elements (Nd, Gd, Pu, etc.), can be achieved. Additionally, information on the distribution of isotopes can then be used for the investigation of fuel burn-up rates for fuel elements placed at different rod positions in the reactor core. Large differences in transmission spectra allow very accurate mapping of isotopic distributions in the samples using either transmission radiography or neutron resonance absorption characteristics of the respective isotopes (Tremsin et al., 2013). One of the attractive features of energy-
Figure 7 – Grain size distribution of uraninite in the matrix and the carbonaceous matter (Sebola, 2014) VOLUME 115
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concentration of uranium in the ore can range from 0.03 to 20%. Conventional mining is by open cut or underground methods. Uranium ore can be produced from a mine specifically for uranium, or as a by-product from mines with a different main product such as copper, phosphate, or gold (International Atomic Energy Agency, n.d.). Using micro-focus X-ray CT with 100 kV potential to distinguish between gold (μρX = 358 cm-1) and uraninite (μρX = 283 cm-1), both the minerals are easy to distinguish from the matrix minerals such as pyrite (μρX = 18.4 cm-1), zircon (μρX = 45.9 cm-1), and brannerite (μρX = 89.6 cm-1) due to their much higher elemental densities. The use of CT as a sorting method is still a challenge, as it is difficult, at the 100 kV X-ray energy CT capability available, to distinguish between gold and uraninite (Chetty et al., 2011). In a followup 3D micro-focus X-ray computer tomography (μXCT) study using 120 kV, the contrast and resolution of the minerals were well defined and individual minerals could be separated and distinguished from other minerals (Sebola, 2014). For the detection of uranium, 3D-CT was benchmarked against 2D mineralogical results from optical microscopy and scanning electron microscopy (SEM). Uraninite, brannerite, and uraniferous leucoxene are the uranium-bearing minerals present in the samples (from the Vaal Reef) and were quantified by μXCT-3D analysis for their sizes, shapes, and distribution with respect to other mineral components in the samples. Uraninite was found to be the major mineral, occurring mainly in the quartz matrices and also associated with carbonaceous matter as depicted in Figure 6. The uraninite and gold in the matrix occurred as rounded grains up to 200 μm in size, as observed by 2D mineralogical techniques. CT allows for 3D grain size analyses of the uraninite grains in the total volume of the sample, and in this study also their association with matrix minerals, as depicted in Figure 7. The CT observations supported the results acquired by conventional mineralogical techniques, suggesting that 3D μXCT can be used to complement other mineralogical techniques in obtaining 3D information. However, 3D μXCT has limitations such as spatial resolution, partial volume effect, and overlapping of mineral grey-scale values. It is therefore suggested that the
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Figure 8 – Neutron radiograph (film) of fresh fuel pellets in a fuel pin with varying degrees of isotopic U-235 enrichment. Pellets with higher enrichment appear darker (Frajtag, n.d.)
Figure 9 – Neutron attenuation of 100 μm thick U-235 and U-238 isotopes calculated from the tabulated data on the total cross-sections as a function of neutron energy (Tremsin et al., 2013)
Figure 10 – Thermal neutron transmission radiographs obtained by grouping the energy-resolved images of different neutron spectra. The ranges of neutron energies used to build each radiograph are shown in the respective legends (Tremsin et al., 2013)
resolved neutron radiography is the ability to enhance contrast, and in some cases enable quantification, as shown in Figure 10. It was observed that the contrast between the pellets of different density depends strongly on the range of neutron energies used. The more thermal part of the beam spectrum (neutron energies above 19.7 MeV) reveals the pellet with the lowest density as an object with the highest transmission. The coldest part of the neutron spectrum (neutron energies even below 6 MeV) shows the least dense pellet as a darkest in the assembly.
Nuclear fuel fabrication and testing (I & PIE) Nuclear fuel types range from isotopic sources in a form of salt or disks to pressurized water reactor (PWR) fuel in the form of UO2 fuel pellets inside ZrTM -cladded fuel pins. Nuclear fuel is subjected to stringent manufacturing and
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performance criteria which have to be verified. Nondestructive testing of the fuel ensures that other tests can be performed subsequently and that the material can still be applied in its specific environment. Some of the NDT tests are applied to characterize and/or quantify the integrity of the fuel or as quality assurance tests. X-ray radiography cannot be used for irradiated fuel inspection, whereas neutron radiography becomes possible due to the following reasons (Lehmann, Vontobel, and Hermann, 2003): ➤ Uranium has a very high attenuation coefficient for Xrays (about 50 cm-1 at 150 keV). The diameter of fuel pellets is in the order of 10 mm and penetration by Xrays is impossible. High-energy gamma radiation (>1 MeV) is, however, suitable for quality control of fresh fuel pellets ➤ The neutron attenuation coefficient for the natural The Journal of The Southern African Institute of Mining and Metallurgy
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➤
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composition of uranium is low (0.8 cm-1) and it is easy to transmit neutrons through thicker assemblies U-235 and U-238 have very different interactions with thermal neutron beams. Due to the 60 times higher cross-section of U-235, it is very easy to distinguish between the two isotopes and to quantify the amount of the fissile isotope U-235 Lead is used as shielding material around fuel samples for radiation protection purposes, and thus X-ray radiography fails in transmission experiments. Neutrons, on the other hand, penetrate lead shielding with a thickness of about 15 cm and allow neutron radiography investigations Additional substitutes in fuel compositions, which are in use as burnable poisons but are strong neutron absorbers (e.g. B, Li-6, Dy, or Gd), are easily identified with neutron methods After long-term exposure, hydrogen can be found in the cladding outer region of fuel rods under some circumstances. X-ray radiography fails to visualize these material modifications because of the very low contrasts obtained for elements with with low atomic numbers. Neutrons, on the other hand, have a high sensitivity for hydrogen, thus allowing quantification of the hydrogen content in cladding.
Isotopes Characterization of isotopic sources is a demanding and difficult task due to their physical size and natural radioactivity, which makes visual inspection impossible. Hoffman (2012) investigated a small radioactive radium source using high-resolution micro-focus X-ray tomography to determine whether the sample contained a powdered form of radioactive material or whether it was solid. The source contained 20 mg of radium and was in a form of a needle with a diameter of about 1 mm and about 8 mm in length – all sealed within a glass tube. Figure 11 is an XCT tomogram of the ampoule showing its serial number clearly on the outside, while Figure 12 is a slice from the 3D tomogram revealing the position of the radioactive material inside the needle. Valuable metrology quantitative information could be deducted from the 3D tomogram of the isotope (Table I). The most important aspect for further processing of the isotope is the quantification of the volume of radioactive salt present in the needle.
Table I
Quantitative information deducted from the XCT Description: Information
Value
Unit
Needle inscription Official source activity Contact dose rate Diameter of tube Diameter of internal void Length of tube Length of internal void Wall thickness Length of salt Inner volume Volume of salt
100 20 1.1 1.67 1.03 9.72 7.57 0.34 3.01 6.31 2.51
µg mCi mSv/h mm mm mm mm mm mm mm3 mm3
temperature, pressure, and radiation level. Thermal neutron radiography investigations were conducted with the conventional film technique due to the radioactivity of the objects. The following issues are addressed through the use of neutron radiography: (a) condition of the fuel assembly, including fuel rod condition, (b) detection of leaks such as ingress of water, and (c) quality control, including functional and dimensional evaluation and inspection of irradiation devices and components. Figure 13 shows a neutron radiograph of a fuel pin with pelletized fuel as fabricated (Domanus, 1992). Due to the high radioactivity of nuclear fuel after irradiation, X-ray radiography cannot be used as an investigation technique. Investigations are done within a hot-cell laboratory set-up, which allows for the remote handling of the radioactive fuel (Klopper, De Beer, and Van Greunen, 1998). A research reactor is normally an extension of a hotcell laboratory, as neutron radiography is one of the major analytical probes in the post-irradiation examination (PIE) of nuclear fuel. Typical findings using neutron radiography as an analytical probe on irradiated fuel pins are the condition of the fuel pellets and of the ZrTM -cladding material. Fuel pellet investigations reveal fabricated conditions such as cracks, chips, change of shape or location, voids, inclusions, corrosion, nuclear properties, and coolant. Figure 14 shows examples of these findings on film neutron radiographs.
Nuclear fuel A major field of neutron radiography application is the inspection of nuclear fuel and control rods, reactor materials and components, and of irradiation devices for the testing of nuclear fuels and materials. The fuel rods are used under extreme conditions such as very high power density,
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Figure 11 – μXCT of a Ra isotopic source ampule (Hoffman, 2012)
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Figure 13 – Neutron radiograph of nuclear fuel prior to irradiation (Domanus, 1992)
Figure 14 – Neutron radiographs (film prints) of irradiated nuclear fuel and their conditions (Domanus, 1992). Top: random cracks in pellets, middle: typical longitudinal cracks in pellets
Domanus (1992) describes a number of other fuel pellet properties revealed by neutron radiography, including central voids and the accumulation of Pu in the central void. Fuel rod inspections include deformed cladding, hydrides in cladding, plenum and spring, dislocated disks, condition of the bottom plug, and a picture of a melted thermocouple inside the fuel rod. Lehmann, Vontobel, and Hermann (2003) report on the extensive utilization of the NEUTRA and ICON neutron radiography (Nrad) facilities at Paul Sherrer Institute in Switzerland, where a dedicated detection station is available for the inspection of irradiated fuel assemblies. Aspects such as fuel enrichment, fuel poisoning, and hydrogen content in the fuel cladding are being addressed and investigated by neutron radiography. Due to the importance of fuel cladding investigations, the utilization and function of neutron radiography is addressed in the following paragraph.
Hydrogen embrittlement It is well known that hydrogen agglomeration is deleterious in any material. More than a few hundred ppm hydrogen in the cladding surface of fuel rods at compromises the structural stability of the cladding tube significantly, with the consequence of possible failure, especially when mechanical loading is also involved. The ability of neutrons to penetrate uranium is considerably higher than for X-rays and allows for the structures of the nuclear fuel rods to be inspected. Furthermore, the probability of neutron interaction with hydrogen is very high, while for X-rays it is effectively zero. This allows neutron radiography to be effectively utilized for
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the study of even small quantities of hydrogen ingress in the cladding, which is an important mechanism for cladding embrittlement, as depicted in Figure 15. Furthermore, through proper characterization and with the aid of digital radiographs, Nrad allows for the investigation of the absolute hydrogen content and its distribution. in situ investigations provide new information about the kinetics of hydrogen uptake during steam oxidation and of hydrogen diffusion in zirconium alloys. Nrad-studies are the only way to investigate and understand the phenomenon of hydrogen ingress in the ZrTM cladding. A linear dependence of the total macroscopic neutron cross-section on the H/ZrTM atomic ratio, as well as on the oxygen concentration, was found, while no significant temperature dependence of the total macroscopic neutron cross-sections of hydrogen and oxygen was found, depending that zirconium and oxygen do not change their structures. Additionally, it was found that rapid hydrogen absorption takes place in the absence of the oxide layer covering the metallic surface of the ZrTM cladding (Grosse et al., 2011). Figure 16 displays the results of in-situ Nrad investigations of hydrogen uptake during steam oxidation with the time dependence of hydrogen concentration of ZrTM-4 materials at 1273 K and higher, where a very rapid hydrogen uptake was found in the first couple of seconds after the steam flow was switched on. At temperatures of about 1273 K a phase transformation occurs and is accompanied by a volume change and the formation of a pronounced crack structure. When the cracks are formed, the The Journal of The Southern African Institute of Mining and Metallurgy
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Figure 15 – Neutron radiograph of nuclear fuel and cladding material showing (black spots) hydrogen accumulation within the ZrTM tubing (Frajtag, n.d.)
Figure 16 – Neutron radiography (Nrad) investigation: kinetics of hydrogen uptake and release during steam oxidation (Frajtag, n.d.)
hydrogen uptake increases by nearly an order of magnitude (Grosse et al., 2008; Grosse, 2010). The decrease in hydrogen concentration is due to the consumption of the ß-ZrTM phase, which contains most of the absorbed hydrogen.
Pebble bed modular reactor (PBMR) fuel Figure 17 shows the composition of a 60 mm outer diameter high-temperature reactor (HTR) fuel pebble consisting of thousands of 0.5 mm diameter low-enriched uranium oxide fuel particles with a tri-structural isotropic (TRISO) coating, embedded in a graphitic matrix. The pebble was analysed using X-ray tomography technique prior to irradiation at the SANRAD facility located at the SAFARI-1 nuclear research reactor in South Africa. The aim of the investigation was to observe the homogeneity of the TRISO particles within the carbon matrix and to direct the manufacturing process to ensure the centralization of the fuel within the carbon matrix
(Necsa, 2006). Figure 18 shows the misalignment of the fuel within the carbon matrix of the fuel pebble as well as the location and identification of a TRISO particle within the fuelfree zone. The inhomogeneous distribution of the TRISO particles at the top of the fuel pebble can be clearly seen. Three-dimensional quantitative data of the misalignment of the fuel particles becomes available in the tomograms and is presented in Figure 19, showing the extent of correction in X-, Y- and Z-directions to be introduced in the manufacturing process of the fuel pebble. Lehmann, Vontobel, and Hermann (2003) reported the successful application of neutron tomography to the 3D scanning of PBMR fuel pebbles at the NEUTRA facility of the SINQ spallation source at PSI in Switzerland (see Figure 20). A sphere-type fuel element from the high-temperature reactor (HTR) programme was studied with neutron tomography. This sample is 6 cm in diameter and contains about 8500 individual fuel pebbles (diameter 0.5 mm). No shielding of the fresh fuel element was necessary for the tomographic inspection. The investigation was aimed at the visualization of the 3D distribution of the fuel particles in the graphite matrix in order to determine its uniformity and the fuel sphere’s content of fissile material. TRISO fuel particles are an integral part of the fuel design for current and future HTRs. A TRISO particle comprises four concentric spherical layers encasing a fuel kernel, namely the buffer (porous carbon), inner pyrolytic carbon (IPyC), silicon carbide (SiC), and outer pyrolytic carbon (OPYC) layers (see Figure 17). Each layer performs specific functions. The fuel kernel, consisting of uranium or uranium carbide, provides the fissile material and retains some of the fission products. The buffer layer, a highly porous carbon structure, provides some free volume for gaseous fission products, and protects the SiC layer from damage by high-energy fission products. The IPyC layer provides structural support for the subsequent SiC layer and prevents the chlorine compounds required for SiC deposition interacting with the fuel kernel. The SiC layer forms
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Figure 17 – Composition of PBMR fuel (Weil, 2001)
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Figure 18 – X-ray tomography of a PBMR fuel pebble. Left: the noncentralized fuel sphere within the carbon matrix. Right: Location and identification of a TRISO particle inside the fuel-free zone (Necsa, 2006)
Figure 20 – Neutron tomogram generated at PSI, Switzerland showing the exact location and homogeneity of the TRISO (Lehmann, Vontobel, and Hermann, 2003)
Figure 19 – Graphical presentation of the deviation of the fuel zone of a BPMR pebble from the centre of the carbon matrix in three dimensions (Necsa, 2006)
the main diffusion barrier for fission products. It acts as a pressure vessel, providing mechanical strength for the particle during manufacture of the nuclear fuel compact or pebble bed. The OPyC layer protects the SiC layer during fuel fabrication as the TRISO particle is pressed into a larger fuel compact or pebble. Lowe et al., (2015) examined the applicability of multiscale X-ray computer tomography (CT) for the non-destructive quantification of porosity and thickness of the various layers of TRISO particles (see Figure 21) in three dimensions, and compared this to the current destructive method involving high-resolution SEM imaging of prepared cross-sections. An understanding of the thermal performance and mechanical properties of TRISO fuel requires a detailed knowledge of pore sizes, their distribution, and interconnectivity. Pore size quantification (false color coding) and distribution in an X-ray tomogram of the SiC (D) and OPyC (E) layers within a TRISO particle is shown in Figure 22. Direct comparison with SEM sections indicates that destructive sectioning can introduce significant levels of coarse damage, especially in the pyrolytic carbon layers. Since it is non-destructive, multi-scale time-lapse X-ray CT opens the possibility of intermittently tracking the degradation of TRISO structure under thermal cycles or radiation conditions in order to validate models of degradation such as kernel movement. Xray CT in situ experimentation on TRISO particles under load and temperature could also be used to understand the internal changes that occur in the particles under accident conditions.
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Figure 21 – Micro-focus X-ray radiograph of a TRISO particle (Necsa)
Figure 22 – Xradia VersaXRM orthoslice showing the SiC (D) and OPyC (E) layers with the pore volumes superimposed. Pore diameter is colourcoded to identify large pore clustering (Lowe et al., 2015)
Research reactor control rod verification Nrad is being applied as a verification and analytic technique at the SAFARI-1 nuclear research reactor on the control rods The Journal of The Southern African Institute of Mining and Metallurgy
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Radioactive waste Low- and intermediate-level nuclear waste is normally encapsulated in in some form of barrier to protect the waste from the environment and vice versa. Intermediate-level nuclear waste is firstly encapsulated in a steel drum, compressed, and finally embedded normally in a concrete drum and safely stored underground in a remote location such as Vaalputs in the Karoo region in South Africa (Necsa, n.d. (b)) (see Figure 24). A site is normally chosen with low rainfall and suitable surface and groundwater conditions. Concrete is a porous medium and the characterization of transport of water through concrete structures is well described by De Beer, Strydom and Griesel (2004) and De Beer, Le Roux and Kearsley (2005). It is especially important to understand the transport of water through concrete because nearly all concrete structures contain steel reinforcing, and in the case of nuclear waste, intermediate-level nuclear waste in compressed steel drums. When cracks in the concrete, caused by the transport of liquid through it, reach the reinforcing, an environment conductive to the corrosion of steel is created. Corrosion affects the strength of the structural members, as the steel is a major contributor to the tensile and compressive strength of the members. Severe leakage of radioactive materials into the surrounding environment is thus possible if the integrity of the concrete barrier is compromised. Neutron radiography studies of concrete and mortars enable the direct physical visualization and quantitative detection of water inside concrete structures. The physical
Figure 23 – Neutron radiographs of a Cd standard (left) and control rod (right) showing Cd (black) indicating high thermal neutron absorbing materials (Necsa, 2010 (a)) The Journal of The Southern African Institute of Mining and Metallurgy
properties of concrete such as porosity, permeability, and sorping characteristics are obtained through applying neutron radiography as a non-destructive analytic tool. The aim of these investigations is to maximize the properties to prevent water sorption and leaching of concrete structures and optimize one of the physical properties which is sometimes neglected in the criteria to develop structures for nuclear waste encapsulation (De Beer, Strydom and Griesel, 2004 ; De Beer, Le Roux and Kearsley, 2005). To improve the durability of concrete, the capillary and pore size within the concrete matrix must be restricted to a minimum. This is why hydration as well as W/C ratio properties is of great importance, and creates thus an ideal opportunity for neutron radiography to play a role in obtaining the needed information in a non-destructive manner to optimize these parameters. The visualization of the sorption of water by means of neutron tomography of a laboratory-size concrete structure is depicted in Figure 25.
Conclusions X-ray and neutron radiography in two or three dimensions play an important role in many dedicated areas within the nuclear fuel cycle. The advantage of these methods is their completely non-destructive nature. Visualization of the structure of samples, as well as quantitative description, are important aspects in materials research. The important roles of X-ray and neutron radiography/tomography as non-invasive analytic techniques within specific areas within the nuclear fuel
Figure 24 – Vaalputs intermediate-level waste storage site, South Africa (Necsa)
Figure 25 – Neutron radiographs showing the effect of 70%, 60%, and 50% W/C ratio on the sorptivity of water into a concrete slab (De Beer, Strydom and Griesel, 2004) VOLUME 115
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prior to their installation in the core of the reactor. The quality control assurance test entails the verification of the neutron attenuation cross-section of the control rod against a standard consisting of Cd. The inspection entails a visual clarification of the attenuation of the thermal neutrons by inspection of the neutron radiograph of the control rod. Additionally, due to the digital radiography capability of the neutron camera detection system, the first-order neutron transmission calculation can be made using the pixel greyscale values on the radiographs of both the standard and control rod sample. Pixel greyscale values represent a linear relationship in the neutron attenuation of materials. In this instance a dramatic decline in greyscale pixel values is seen due to the high thermal neutron absorption by the Cd section (μρn =115.11cm-1) of the control rod. Neutron radiographs of the Cd standard and a control rod are depicted in Figure 23.
Neutron- and X-ray radiography/tomography cycle should not be underestimated. X-rays and neutrons are produced by very different methods, and also interact with materials in different manners. In the nuclear environment, each type of radiation has its own field of utilization due to their different characteristics, but in some instances their applications complement each other to reveal comprehensive information. Neutron transmission analysis is a very helpful tool to obtain information on the properties of, and changes in, nuclear fuel material. Scientists and researchers in the geosciences in South Africa have, in the availability of the tomography facilities at Necsa, the capabilities to conduct quantitative analytical measurements at state-of-the-art radiation imaging facilities that compare to similar facilities elsewhere in the world. Within the mining area, 3D computer tomography shows potential for further development, and can be already used to complement and add value to current conventional 2D mineralogical techniques. Neutron radiography analysis is able to derive the hydrogen content in fuel cladding both qualitatively and quantitatively, with high sensitivity and precision. The results presented here illustrate how recent advances in laboratory-based X-ray CT instruments allow the examination of TRISO particles at the nano- and micro-scales in 3D. In this case study, high-resolution X-ray CT has been shown to be a viable tool for profiling the TRISO particles in two important aspects; to characterize the individual TRISO layers with variations in thickness and their subsequent interactions, thus allowing manufacturing validation as well as assisting in working towards a mechanistic understanding of fabrication and in-service issues. The availability of these techniques in South Africa opens new possibilities for research, quantitative analysis, and nondestructive evaluation. National capacity as well as international trends shows the ability for non-destructive testing of nuclear materials utilizing penetrating X- ray- and neutron radiation in more comprehensive and unique ways than before.
References AFRICAN NTD CENTRE. http://www.andtc.com/ BANHARD, J. 2008. Advanced Tomographic Methods in Materials Research and Engineering. Oxford University Press. CHETTY, D., CLARK, W., BUSHELL, C., SEBOLA, P.T., HOFFMAN, J.W., NSHIMIRIMANA, R.B., and DE BEER, F.C. 2011. The use of 3D X-ray computed tomography for gold location in exploration drill cores. Proceedings of the 10th International Congress for Applied Mineralogy (ICAM), Trondheim, Norway, 1-5 August 2011. pp.129–136. http://www.mintek.co.za/wpcontent/uploads/2011/11/ch-17.pdf DAVIES, L.M. 2000. Role of NDT in condition based maintenance of nuclear power plant components. 15th World Conference on Non-Destructive Testing, Rome, 15-21 October 2000. http://www.ndt.net/article/wcndt00/papers/idn078/idn078.htm DE BEER, F.C. and AMEGLIO, L. 2011. Neutron, X-ray and dual gamma-ray radiography and tomography of geomaterial – a South African perspective. Leading Edge, June 2011. Special Edition Africa. DE BEER, F.C., MIDDLETON, M.F., and HILSON, J. 2004. Neutron radiography of porous rocks and iron ore. Applied Radiation and Isotopes, vol. 61. pp. 487–495. DE BEER, F.C., STRYDOM, W.J., and GRIESEL, E.J. 2004. The drying process of concrete: a neutron radiography study. Applied Radiation and Isotopes, vol. 61, no. 4. pp. 617–623. DE BEER, F.C., LE ROUX, J.J., and KEARSLEY, E.P. 2005. Testing the durability of concrete with neutron radiography. Nuclear Instruments and Methods in Physics Research A, vol. 542. pp. 226– 231. DOMANUS, J.C. 1992. Practical Neutron Radiography. Kluwer Academic Publishers. FRAJTAG, P. Not dated. Radiation Protection and Radiation Applications: Gamma and Neutron Radiography. http://moodle.epfl.ch/pluginfile.php/1593971/mod_resource/content/2/RRA -EPFL- FS2014-Week14a.pdf [Accessed 1 May 2015].
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GHOSH, J.K., PANAKKAL, J.P., and ROY, P.R., 1983. Monitoring plutonium enrichment in mixed-oxide fuel pellets inside sealed nuclear fuel pins by neutron radiography. NDT International, vol. 16, no. 5, October 1983. pp. 275–276. DOI: 10.1016/0308- 9126(83)90127-X GROSSE, M. 2010. Neutron radiography—a powerful tool for fast, quantitative and non- destructive determination of the hydrogen concentration and distribution in zirconium alloys. Proceedings of the 6th International ASTM Symposium on zirconium in the Nuclear Industry, Chengdu, China, 9–13 May 2010. GROSSE, M., VAN DEN BERG, M., GOULET, C., LEHMANN, E., and SCHILLINGER, B. 2011. in situ neutron radiography investigations of hydrogen diffusion and absorption in zirconium alloys. Nuclear Instruments and Methods in Physics Research, vol. A651. pp. 253–257. doi: 10.1016/j.nima.2010.12.070 GROSSE, M., STEINBRUECK, M., LEHMANN, E., and VONTOBEL, P. 2008. Kinetics of hydrogen absorption and release in zirconium alloys during steam oxidation. Oxidation of Metals, vol. 70, no. 3. pp. 149–162. GRÜNAUER, F. 2005. Design, optimization, and implementation of the new neutron radiography facility at FRM-II. Dr. Rer. Nat. dissertation, Faculty of Physics, Technischen Universität München. HOFFMAN, J.W. 2012. Process description for micro-focus X-ray investigation of source at MIXRAD facility. Necsa internal report, DOC NO: RS-MFX-PRO12002. INTERNATIONAL ATOMIC ENERGY AGENCY. Not dated. Getting to the Core of the Nuclear Fuel Cycle. Department of Nuclear Energy, Vienna, Austria. https://www.iaea.org/OurWork/ST/NE/NEFW/_nefwdocuments/NuclearFuelCycle.pdf ISO 15708-1:2002(E). Non-destructive testing — Radiation methods — Computed tomography — Part 1: Principles. KLOPPER, W., DE BEER, F.C., and VAN GREUNEN, W.S.P. 1998. Overview of hot cell facilities in South Africa. Proceedings of the European Working Group ‘Hot Laboratories and Remote Handling’. LEHMANN, E.H., VONTOBEL, P., and HERMANN, A. 2003. Non-destructive analysis of nuclear fuel by means of thermal and cold neutrons. Nuclear Instruments and Methods in Physics Research A, vol. 515, no. 3. pp. 745–759. doi:10.1016/j.nima.2003.07.059 LOWE, T., BRADLEY, R.S., YUE, S., BARII, K., GELB, J., ROHBECK, N., TURNER, J., and WITHERS, P.J. 2015. Microstructural analysis of TRISO particles using multiscale X-ray computed tomography. Journal of Nuclear Materials, vol. 461. pp. 29–36. http://dx.doi.org/10.1016/j.jnucmat.2015.02.034 MCGLINN, P.J., DE BEER, F.C., ALDRIDGE, L.P., RADEBE, M.J., NSHIMIRIMANA, R.B., BREW, D.R.M., PAYNE, T.E., and OLUFSON, K.P. 2010. Appraisal of a cementitious material for waste disposal: neutron imaging studies of pore structure and sorptivity. Cement and Concrete Research, vol. 40. pp. 1320–1326 NECSA. 2006. RT-TVG-06/05: Test Report: X-Ray Radiography of PBMR-Fuel Spheres with Zirconium Particles. Internal report, SAMPLE NO: DFS-TF03G04. NECSA. 2010 (a). RS-TECH-REP-10004: Neutron Radiography Quality Assurance test report of neutron absorbing material in Control Rods of the SAFARI-1 Nuclear Research Reactor (Lot/Batch No: RT-LOT-10/03). Internal report. NECSA. Not dated (b). Vaalputs. The National Radioactive Waste Disposal Facility. http://www.radwaste.co.za/vaalputs%20information%20pamflet.pdf. [Accessed 1 May 2015]. OECD NUCLEAR ENERGY AGENCY. 2003. Nuclear Energy Today. OECD Publishing. p. 25. PIXSHARK. Not dated. http://pixshark.com/images-of-nuclear-fuels.htm SAIW (Southern African Institute of Welding). Not dated. http://www.saiw.co.za/ SEBOLA, P. 2014. Characterisation of uranium-mineral-bearing samples in the Vaal Reef of the Klerksdorp Goldfield, Witwatersrand basin. MSc dissertation, Faculty of Science, University of the Witwatersrand, Johannesburg. http://wiredspace.wits.ac.za/handle/10539/16820?show=full http://hdl.handle.net/10539/16820 SGS SOUTH AFRICA (PTY) LTD. Not dated. http://www.sgs.co.za/en.aspx TREMSIN, A.S., VOGEL, S.C., MOCKO, M., BOURKE, M.A.M., YUAN, V., NELSON, R.O., BROWN, D.W., and FELLER, B. 2013. Non-destructive studies of fuel pellets by neutron resonance absorption radiography and thermal neutron radiography. Journal of Nuclear Materials, vol. 440. pp. 633–646. WEIL, J. 2001. Pebble-bed design returns. Nuclear Power gets a Second Look. IEEE Spectrum Special Report. http://spectrum.ieee.org/energy/nuclear/ pebblebed-design-returns [Accessed 20 June 2015]. WILLCOX, M. and DOWNES, G. Not dated. A brief description of NDT techniques. Insight NDT, paper T001. http://www.turkndt.org/sub/makale/ornek/a%20brief%20description%20of %20ndt.pdf WORLD NUCLEAR ASSOCIATION. 2015. http://www.world- nuclear.org/info/NuclearFuel-Cycle/Mining-of-Uranium/World-Uranium-Mining- Production/ [Accessed 13 May 2015]. ◆ The Journal of The Southern African Institute of Mining and Metallurgy
http://dx.doi.org/10.17159/2411-9717/2015/v115n10a4
Non-destructive characterization of materials and components with neutron and X-ray diffraction methods by A.M. Venter*
Background The availability of advanced characterization techniques is integral to the development of advanced materials, not only during development phases, but in the manufactured components as well. At Necsa, two modern neutron diffractometers equipped with in-situ sample environments, as well as complementary X-ray diffraction instruments, are now available as User Facilities within the National System of Innovation in support of the South African research and industrial communities. Neutrons and X-rays, owing to their different interaction mechanisms with matter, offer complementary techniques for probing crystalline materials. Both techniques enable nondestructive investigation of phenomena such as chemical phase composition, residual stress, and texture (preferred crystallite orientation). More specifically, the superior penetration capabilities of thermal neutrons into most materials allows for the analysis of bulk or localized depth-resolved properties in a wide variety of materials and components. Materials that can be investigated include metals, alloys, composites, ceramics, and coated systems. In particular, depth-resolved analyses using neutron diffraction complements surface investigations using laboratory X-rays in many scientific and engineering topics. The diffraction techniques can add significant downstream value to the anticipated nuclear industry development activities. Keywords residual stress; crystallographic texture; chemical phase identification, neutron and X-ray diffraction.
Introduction Materials characterization is central in understanding the relationship between the structure, properties, and performance in order to engineer materials that fit the performance criteria for specific applications. This is conveniently represented in the form of a tetrahedron, generally known as the material science paradigm (Figure 1). The availability of advanced characterization techniques is integral to the development of advanced metals, not only during development phases, but in the form of manufactured components. At Necsa (South African Nuclear Energy Corporation), modern X-ray and neutron diffraction instruments are now accessible to the South African research and industrial communities. These facilities enable nondestructive investigations of materials and components that could add significant downstream value to the anticipated nuclear industry development activities. The Journal of The Southern African Institute of Mining and Metallurgy
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Approximately 95% of solid materials can be described as crystalline. By applying the wave properties of X-rays (keV energies as generated in laboratory-based units, or highpower synchrotron facilities) and neutrons (meV energies) having wavelengths of the same order as atomic spacings, i.e. 10-10 m, unique information can be obtained. Their interaction with the crystalline ordering leads to constructive interference described by Bragg’s law of diffraction. This leads to characteristic diffraction patterns for each phase that essentially are fingerprints of the chemical phase content (Fourier transform of the atomic arrangement). Both the peak positions (corresponding to lattice plane spacings) and the relative intensities (corresponding to the atomic species and arrangement in the unit cell) in the diffraction patterns are indicative of specific phases. Xray photons scatter through an electromagnetic interaction with the electron charge cloud of the material, while neutrons are scattered by interaction with the nuclei. In general the interaction strength of X-rays with matter is directly related to the atomic number of the materials being investigated, whereas neutron scattering lengths are approximately equal in magnitude for most atoms. Apart from the interaction strength differences, their penetration depths are also dependent on the atomic species, as summarized in Table I for a number of technologically important materials. The different interaction mechanisms of neutrons and X-rays with matter offer complementary techniques for investigating crystalline materials at the microstructural
* Research and Development Division, The South African Nuclear Energy Corporation SOC Ltd., (Necsa), Pretoria, South Africa. Š The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. OCTOBER 2015
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Synopsis
Non-destructive characterization of materials and components
Figure 1 – Material science paradigm
level. Both techniques enable nondestructive investigation of phenomena such as chemical phase composition, residual stress, and texture (preferred crystallite orientation). Materials include metals, alloys, composites, ceramics, and coated systems. In particular, depth-resolved analysis using neutron diffraction complements surface investigations using laboratory X-rays in many scientific and engineering disciplines. Neutrons complement X-rays due to the following properties: ➤ Neutrons are electrically neutral. This enables orders of magnitude deeper penetration of bulk material/components: • Allows nondestructive bulk analysis • Ease of in situ experiments, e.g. variable temperature, pressure, magnetic field, chemical reaction, etc. ➤ Neutrons detect light atoms even in the presence of heavy atoms (organic crystallography) – especially hydrogen. This property has been decisive in the investigation of high-temperature superconductors ➤ Neutrons distinguish atoms adjacent in the periodic table, and even isotopes of the same element (changing scattering picture without changing chemistry). This is particularly applicable to the transition metal series ➤ Neutrons have a magnetic moment. This enables nondestructive investigation of magnetic phenomena from direct observation of the reciprocal lattice.
Neutrons of adequate flux for neutron diffraction applications are produced as a by-product of the fission of 235U in neutron research reactors (such as SAFARI-1 operated by Necsa in South Africa), or in accelerator-based facilities when very high-energy protons strike a target producing ‘spallation’. These high-energy (MeV) neutron products are then thermalized and filtered to the thermal energy range (meV). The neutron diffraction instruments at SAFARI-1 operate in constant wavelength mode where a highly monochromatic thermal neutron beam (<1% wavelength spread with wavelengths selectable in the range 1.0–2.0 Å, typically 25 meV energies) with a flux in the order of 106 neutrons per square centimetre per second is extracted from the fission energy spectrum and directed to a sample. Bragg’s law of elastic diffraction nλ = 2dhklsinθhkl describes the geometrical condition for coherent diffraction that can be measured to high precision on a diffractometer. In this equation λ is the monochromatic wavelength in Å (10-10 m), dhkl is the interatomic spacing between the parallel crystallite planes of Å dimension (hkl refers to the Miller notation of the crystal planes) and θhkl is the angle at which the diffracted peak is measured. In the angular dispersive operational mode of the SAFARI-1 neutron diffraction instruments, the wavelength is selected with the instrumental setup and thus accurately known, with dhkl and θhk being the only variables. The latter is measured experimentally to high precision with the diffraction instrument, which comprises a goniometer for sample positioning, and a detector that is precisely rotated around the goniometer axis in the horizontal plane. The diffracted intensity is measured with a high-sensitivity neutron detector that uses 3He as ionization medium for the accurate measurement of the diffraction angles from all coherently scattered Bragg peaks from which respective values can be calculated. As is the case with X-rays, all crystalline materials (even chemically multi-phased materials) placed in the neutron beam produces a diffraction pattern. This document provides an overview of a subset of diffraction techniques that may be of relevance to the materials beneficiation aims of the Advanced Materials Initiative.
Table I
Neutron and X-ray scattering parameters for a number of elements that comprise the common engineering alloys. Half attenuation lengths correspond to the thickness of material that reduces the intensity by 50%. 1.5 Å X-rays correspond to Cu radiation (8 keV). 0.15 Å X- rays correspond to synchrotron radiation (100 keV) Element
Neutron coherent scattering length (bc)
1.8 Å neutrons [mm]
Half attenuation lengths 1.5 Å X-rays [μm]
0.15 Å X-rays [mm ]
[fm = 10-15 m ]
X-ray scattering length (f0 at sin θ/λ = 0.2 [fm = 10-15 m]
5.375 3.449 -3.37 9.54 2.49 10.3 7.16 7.77 4.86 8.417
0.246 0.258 0.451 0.565 0.594 0.624 0.884 1.715 1.762 2.172
43 66 12 20 2 3 24 1 5 9
100 52 11 20 14 16 8 3 2 1
24 15 6 2.4 2 1.8 1.1 0.1 0.1 0.2
Mg Al Ti Fe Co Ni Zr Hf W U
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Non-destructive characterization of materials and components Applications Results from selected typical applications pursued at the Necsa facilities are reported to provide an indication of the potential value adding that these techniques could offer.
Chemical phase identification The most widespread use of powder (polycrystalline) diffraction is in chemical phase analysis. This encompasses phase identification (search/match), investigation of highand low-temperature phases, solid solutions, and determinations of unit cell parameters of new materials. A multiphase mixture, e.g. a soil sample, will show more than one pattern superposed, allowing for determination of the relative concentrations of phases in the mixture. The powder diffraction method is well suited for characterization and identification of polycrystalline phases. This is done against the International Center Diffraction Data (ICDD) database, which contains in excess of 50 000 inorganic and 25 000 organic phases. Figure 2 shows a neutron diffraction pattern of a multi-phased sintered iron powder sample that was investigated to quantify the constituent phases. The quantification was done with the Rietveld profile refinement technique (Rietveld, 1969; Taylor, 2001), in which a theoretical line profile (a combination of the crystal system and instrumental model) is matched in a least-squares refinement approach to the measured data. Such quantitative phase analysis is now routinely used in industries ranging from cement manufacture to the oil industry and can provide detection limits of 1 weight per cent (wt.%).
Residual stress The total stress that a component experiences in practical use is the vector sum of the applied and residual stresses. The applied stresses result from the loading forces in use and can be calculated to high precision. Residual stresses, consisting of locked-in stress that remains in a material or component after the external forces that caused the stress have been removed, are mostly only approximated qualitatively. These stresses can be introduced by any mechanical, chemical, or thermal process, such as machining, plating, and welding. Stress analysis by diffraction techniques is based on accurate measurement of lattice strain distributions (variations in interatomic spacing). Using Hooke’s law, stresses are
calculated from the strain distributions. Residual stresses are a double-edged sword in material science applications. Compressive residual stresses are beneficial in applications where fatigue performance is required, being able to mitigate crack initiation and propagation. Tensile residual stresses are generally considered to be detrimental as they can lead to crack initiation and propagation. Specifically, residual stress tailoring can render substantial improvements in the optimization of component design. X-ray and neutron radiation enable nondestructive probing at different penetration depths (Hutchings et al., 2005; Fitzpatrick and Lodini, 2003; Reimers et al., 2008; Webster, 2000; Ohms et al., 2008; Engler and Randle, 2010) (Table I). The stress tensors are obtained from the measured strain tensors and application of Hooke’s law of elasticity
where S1 and 1/2S2 are the diffraction elastic constants. Examples from recent investigations performed using neutrons are presented to illustrate potential applications.
Welded mild steel plate As part of an investigation into the influence of differential hardness between the weld metal and base metal on residual stress and susceptibility to stress corrosion cracking, a comprehensive two-dimensional (2D) mapping of the residual stress field, through the 17 mm plate thickness and across the weld was completed. Figure 3a indicates the 2D map of the longitudinal stress component, which is maximally influenced due the differential longitudinal contraction. Values can be as large as the yield stress of the material. Figure 3b indicates the 2D map of this stress component after the sample had undergone a post-weld heat treatment. This treatment has been very successful in having completely relaxed the tensile stress values. Such analyses are only possible owing to the penetrating capabilities of thermal neutrons and the nondestructive nature of the measurement technique.
Laser shock-peened aluminium plate Laser shock-peen (LSP) is an emerging cold work process used to induce compressive residual stresses in metallic
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Figure 2 – Neutron powder diffraction results (measured diffracted intensities as function of diffraction angle) of a mixed-phase sintered iron sample, shown as dots (underneath the red solid line at positive intensities). The multiphase quantification results from a Rietveld analytical approach (red solid line at positive intensities) are indicated in the legend. The curve along the zero intensity level depicts the difference between the measured and analysed data and gives the goodness-of-fit. Vertical lines at negative intensities represent the Bragg peaks (hkl) corresponding to each phase
Non-destructive characterization of materials and components
Figure 3 – Two-dimensional maps of the longitudinal residual stress component in SA 516 Gr 70 pressure vessel steel welded plates: (a) as-welded condition, (b) post-weld heat-treated condition. The weld centre-line is at 0 mm
Figure 4 – Through-thickness stress profile in aluminium plate: (a) parent material, (b) laser shock-peened plate, (c) net contribution by laser shock-peen
Figure 5 – 2D map of the longitudinal residual stress component in aluminium laser beam welded plates: (a) as-welded condition, (b) post-weld-treated condition after laser shock-peen treatment. The weld centre-line is at 0 mm
components to depths up to 1 mm. The purpose of this treatment is to reduce or remove tensile stresses in the surface region to improve the fatigue performance. Figure 4 summarizes the through-thickness in-plane residual stress distribution in an as-rolled aluminium AA6056-T4 plate, taken as the reference, and the stress distribution after laser shock-peen using a 3 GW laser beam. The net effect, purely due to the peen action, was determined by subtracting the reference values.
Laser-welded aluminium with laser shock peen treatment Laser beam welding is a newly established joining technology using no filler material. Similar to other welding techniques,
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it introduces adverse tensile residual stresses as well as component distortions. Figure 5 indicates 2D maps of the residual stress field measured through a 3.3 mm thick AA6056-T4 aluminium plate and across the weld. Here again, the longitudinal stress component is most dominant, as shown in Figure 5a. To improve this adverse stress condition, the weld region has been subjected to laser shock-peen treatment using a 3 GW laser pulse. Compared to the post-weld heat treatment results given previously, the laser treatment has completely altered the tensile stress so as to be substantially compressive.
Additive-manufactured Ti-6Al-4V In the selective laser beam melt (SLM) additive manufacThe Journal of The Southern African Institute of Mining and Metallurgy
Non-destructive characterization of materials and components
Figure 6 – Measured 2D stress component mapped perpendicular to the build direction in a SLM produces Ti-6Al-4V sample: (a) sample geometry showing the ‘internal plane’ investigated nondestructively, (b) 2D residual stress map
turing process, a high-power laser beam is used to melt and deposit successive layers of powder to form complex threedimensional metal parts. The highly localized heat input leads to large thermal gradients. This in turn produces complex residual stress states inside the part. Figure 6 shows a stress map measured in the centre of the sample perpendicular to the build direction in a Ti-6Al-4V sample. This reveals the existence of tensile residual stresses in the nearsurface regions, while the central region of the sample is in compression.
Texture analysis Most solid-state matter has a polycrystalline structure composed of a multitude of individual crystallites or grains. The crystallographic orientation can have various arrangements, ranging from completely random to the development of preferred alignment. The significance of this texture lies in the corresponding anisotropy of many material properties. The influence could be as much as 20–50% of the property value. X-ray and neutron diffraction analysis methods are well established, rendering results referred to as
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Figure 7 – Pole figure representations of selected reflections from two ferritic stainless steel samples subjected to different heat treatments, measured with neutron diffraction
Non-destructive characterization of materials and components macro or bulk texture (Engler and Randle, 2010). Analyses can also be supplemented by methods whereby individual orientations are measured by transmission or scanning electron microscopy and are directly related to the microstructure, which has given rise to the term ‘microtexture’. The preferred orientation is usually described in terms of pole figures. The inverse pole figure gives the probability of finding a given specimen direction parallel to crystal (unit cell) directions. By collecting data for several reflections and combining several pole figures, the complete orientation distribution function (ODF) of the crystallites within a single polycrystalline phase that makes up the material can be determined (Engler and Randle, 2010). Shown in Figure 7 are bulk pole figures measured with neutron diffraction on ferritic stainless steel specimens that have been subjected to different heat-treatment parameters after rolling. Significant texture development is evident, which will contribute to anisotropic properties that could lead to physical phenomena such as earing, warping, and general dimensional changes and distortions.
REIMERS, W., PYZALLA, A.R., SCHREYER, A., and CLEMENS, H. (eds). 2008. Neutrons and Synchrotron Radiation in Engineering Materials Science. Wiley-Vch Verlag GmbH.
RIETVELD, H.M. 1969. A profile refinement method for nuclear and magnetic structures. Journal of Applied Crystallography, vol. 2, no. 2. pp. 65–71. doi:10.1107/S0021889869006558
TARAN, Y., BALAGUROV, A., SABIROV, B., DAVYDOV, V., and VENTER, A. 2014. Neutron diffraction investigation of residual stresses induced in niobiumsteel bilayer pipe manufactured by explosive welding. Materials Science Forum, vol. 768-769. pp. 697-704. doi:10.4028/www.scientific.net/MSF.768-769.697
TAYLOR, J.C. 2001. Rietveld made easy: a practical guide to the understanding of the method and successful phase quantifications. Sietronics Pty. Ltd., Canberra.
TROIANO, E., UNDERWOOD, J.H., VENTER, A.M., IZZO, J.H., and NORRAY, J.M. 2012.
Conclusions
Finite element model to predict the reverse loading behavior of
These examples serve to demonstrate the capabilities of the X-ray and neutron diffraction techniques available at Necsa for the characterization of new materials and advanced beneficiation techniques to enhance component performance in various applications. The vast field of applications can be reviewed in the extensive literature references provided, as well as recent publications from Necsa (Taran et al., 2014; Troiano et al., 2012; Venter, Luzin, and Hattingh, 2014; Venter et al., 2008, 2013, 2014a, 2014b; Zhang et al., 2008). Due the nondestructive nature of the techniques, samples can be investigated at various stages of their manufacture or utilization lifetimes.
autofrettaged A723 and Hb7 cylinders, Journal of Pressure Vessels and Piping, PVT-11-1204, 041012-1.
VENTER, A.M., LUZIN V., and HATTINGH, D.G. 2014. Residual stresses associated with the production of coiled automotive springs. Material Science Forum, vol. 777. pp. 78–83. doi:10.4028/www.scientific.net/MSF.777.78
VENTER, A.M., OLADIJO, O.P., CORNISH, L.A., and SACKS, N. 2014a. Characterisation of the residual stresses in HVOF WC-Co coatings and Substrates. Material Science Forum, vol. 768-769. pp. 280–285. doi:10.4028/www.scientific.net/MSF.768-769.280
Acknowledgments Specific acknowledgment is attributed to the various coworkers from Necsa and academia and their willingness that the results may be used in this document: Professor Johan de Villiers (University of Pretoria), Deon de Beer (University of Pretoria), Victoria Cain (University of Cape Town), Daniel Glaser (University of the Witwatersrand), and Deon Marais from Necsa.
VENTER, A.M., LUZIN, V., OLADIJO, O.P., CORNISH, L.A., and SACKS, N. 2014b. Study of interactive stresses in thin WC-Co coating of thick mild steel substrate using high-precision neutron diffraction. Materials Science Forum, vol. 772. pp. 161–165. doi:10.4028/www.scientific.net/MSF.772.161
VENTER, A.M., OLADIJO, O.P., LUZIN, V., CORNISH, L.A., and SACKS, N. 2013. Performance characterization of metallic substrates coated by HVOF
References
WC–Co. Thin Solid Films, vol. 549. pp. 330–339.
ENGLER, O. and RANDLE, V. (eds). 2010. Introduction to Texture Analysis, Macrotexture, Microtexture, and Orientation Mapping. CRC Press, Taylor and Francis.
VENTER, A.M., VAN DER WATT, M.W., WIMPORY, R.C., SCHNEIDER, R., MCGRATH, P.J., and TOPIC, M. 2008. Neutron strain investigations of laser bent samples.
FITZPATRICK, M.E. and LODINI, A. 2003. Analysis of Residual Stress by Diffraction using Neutron and Synchrotron Radiation. Taylor and Francis. HUTCHINGS, M.T., WITHERS, P.J., HOLDEN, T.M., and LORENTZEN, T. (eds). 2005 Introduction to the Characterization of Residual Stress by Neutron Diffraction. Taylor and Francis. OHMS, C., MARTINS, R.V., UCA, O, YOUTSOS, A.G., BOUCHARD, P.J., SMITH, M., KEAVEY, M., BATE, S.K., GILLES, P., WIMPORY, R.C., and EDWARDS, L. 2008. Proceedings of 2008 ASME Pressure Vessels and Piping Conference (PVP 2008), Chicago, Illinois, 27–31 July 2008. (PVP2008-61913).
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Materials Science Forum, vol. 571–572. pp. 63-68.
WEBSTER, G.A. (ed.). 2000. Neutron Diffraction Measurements of Residual Stress in a Shrink-fit Ring and Plug. VAMAS Report No. 38.
ZHANG, S.Y., VENTER, A.M., VORSTER, W.J.J., and KORSUNSKY, A.M. 2008. High energy synchrotron X-ray analysis of residual plastic strains induced in shot peened steel plates. Journal of Strain Analysis, vol. 43, no. 4. pp. 229-241.
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The Journal of The Southern African Institute of Mining and Metallurgy
http://dx.doi.org/10.17159/2411-9717/2015/v115n10a5
Fluorine: a key enabling element in the nuclear fuel cycle by P.L. Crouse
Fluorine – in the form of hydrofluoric acid, anhydrous hydrogen fluoride, elemental gaseous fluorine, fluoropolymers, volatile inorganic fluorides, and more – has played, and still plays, a major role in the nuclear industry. In order to enrich uranium, the metal has to be in the gaseous state. While more exotic methods are known, the standard and most cost-competitive way of achieving this is by means of uranium hexafluoride (UF6). This compound sublimates at low temperatures, and the vapour is enriched using centrifugal processes. The industrial preparation of uranium hexafluoride requires both elemental fluorine gas and anhydrous hydrogen fluoride (HF). HF is prepared by the reaction of sulphuric acid with fluorspar (CaF2). Fluorine gas in turn is prepared by the electrolysis of HF. Yellowcake is first converted to uranium tetrafluoride (UF4), using HF, after which the compound is treated with fluorine to yield UF6. After enrichment, the UF6 is reduced to UO2 for use in fuel elements in pellet form. South Africa has the largest reserves of fluorspar internationally, and is the third largest producer after Mexico and China. Fluorine technology has many associated difficulties, because of the reactivity of fluorine and the toxicity of HF. The main barriers to entry into the fluorochemical industry are thus the abilities to produce both HF and F2. Both these substances are produced locally, at the industrial scale, at Pelchem SOC Ltd. Should South Africa contemplate developing its own nuclear fuel cycle as part of the awaited new-build nuclear project, it will be imperative to leverage the existing skills with respect to fluorine technology, resident at both Pelchem and Necsa, for this purpose. This paper summarizes the fluorochemical skills developed locally over the past several decades, and suggests strategies for maintaining the technology base and developing it for the next generation of scientists and engineers. Keywords flourine, hydrogen flouride, nuclear fuel cycle.
Introduction Fluorine chemistry has always had a very close relationship with the nuclear industry. Although the first mention of fluorspar was recorded in the sixteenth century, and the fundamentals had been well-developed by the Second World War, it was only during the Manhattan Project (Banks et al., 1994), when almost unlimited funds were made available for the large-scale industrialization of fluorine and related compounds, that the technology took off. Expertise in both the chemistry and the engineering aspects of fluorine is essential for the design and running of several of the unit processes in the nuclear fuel cycle. The Journal of The Southern African Institute of Mining and Metallurgy
The aim of this paper is to give a brief overview of the role of fluorine in the nuclear fuel cycle, and to outline the current South African fluorochemical capability.
Hydrogen fluoride and elemental fluorine The major barrier to entry into the fluorochemical industry is the ability to manufacture and, in general, to handle fluorine and its precursor, hydrogen fluoride. Both are extraordinarily difficult and dangerous substances with which to work. Hydrogen fluoride (HF) is produced by the reaction of fluorspar (calcium difluoride) with sulphuric acid: [1]
The reaction is endothermic and reactors are generally run at temperatures above 200°C. HF is a clear liquid, with a boiling point of 19.6°C. It readily dissolves in water and, in its aqueous form, is known as hydrofluoric acid. Unlike the other common mineral acids, it is weak acid, thus does not readily deprotonate. Hydrofluoric acid is distinguished by its ability to dissolve glass, and as a consequence cannot be used in ordinary laboratory glassware. Both hydrogen fluoride and hydrofluoric acid are enormously hazardous substances (Bertolini, 1992; Smith, 2004). Upon contact, human skin is not immediately burnt by the action of the hydronium ion; rather, because of the small size of the HF molecule, it diffuses through the skin and precipitates and inactivates biological calcium and magnesium subcutaneously, causing tissue necrosis. The wounds are extremely painful, and difficult to treat. In general the dead flesh has to be
* Fluoro-Materials Group, Department of Chemical Engineering, University of Pretoria. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. VOLUME 115
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Synopsis
Fluorine: a key enabling element in the nuclear fuel cycle surgically removed, and a calcium gluconate solution injected into the tissue underneath the wound to prevent deeper diffusion of the HF. A typical HF wound is shown in Figure 1. HF is also very corrosive, and materials selection is critically important in the development of an HF-related industrial process. Although HF was first synthesized in the eighteenth century, and was known to contain an element, it was only in the late nineteenth century that fluorine was first isolated. This was accomplished by Henri Moissan (Argawal, 2007). Because it is a weak acid, anhydrous HF (AHF) does not conduct electrical current. Moissan’s discovery was that the molten salt KF.xHF does indeed conduct electrical current and can be electrolysed. As a melt, e.g. with x=2, it dissociates and forms the equilibrium: [2] Moissan (Figure 2) was awarded the Nobel Prize in Chemistry in 1906. Since the first isolation of fluorine, the electrolysis process has undergone a few technical changes, but has remained more or less static throughout the past few decades. Comprehensive descriptions of the technology can be found in Slesser and Schram (1951), Rudge (1971), and Shia (2004). Fluorine itself is the most reactive element in the periodic table, and reacts with all other elements, excluding only the
Figure 1 – A hydrogen fluoride wound being treated
two noble gases helium and neon (Cotton et al., 2007). In general the reactions of fluorine are highly exothermic, and because of its reactivity, materials of construction are of critical importance for safe operation. In general, expensive nickel-containing alloys are required. It should be noted that HF sells for US$1–3 per kilogram, while F2 sells for US$15–20 per kilogram. The high cost of fluorine can be attributed to electrical requirements and the high maintenance costs of the electrolysis cells. South Africa is richly endowed with fluorspar (Roskill, 2009). Relative production and reserve figures are given in Table I. At present South Africa is the third largest producer of the ore, and has the largest reserves. China and Mexico, being closer to the larger international markets, are the two top producers.
A brief history of fluorine chemistry and technology (Ameduri, 2011) A timeline for the major discoveries and developments in fluorine chemistry is listed below. ➤ Georg Bauer first describes the use of fluorspar (CaF2) as a flux in 1530 – as a flux aiding the smelting of ores by German miners ➤ Heinrich Schwanhard finds, in 1670, that fluorspar dissolved in acid and the solution could be used to etch glass ➤ From the 1720s, the effect on glass by adding sulphuric acid to fluorspar is studied ➤ Scheele (a Swedish scientist) ‘discovers’ fluoric acid (HF) in 1771 ➤ Several chemists try unsuccessfully to isolate fluorine, and several die of HF poisoning during separation experiments ➤ The French chemist Moissan is the first to isolate elemental fluorine gas. He is awarded the Nobel Prize in 1906 ➤ Swarts discovers the Cl/F exchange chemistry of SbF3 ➤ Midgley discovers Freons in 1928 ➤ In the 1930s General Motors begins using Freon-type fluorocarbons (CFCs) as replacement for hazardous materials e.g. NH3. CFCs also finds use in propellants and fire extinguishers ➤ 1938 Plunkett of DuPont discovers Teflon® ➤ WW II and uranium enrichment ➤ In 1947 Fowler discovers the CoF3 method of perfluorination
Figure 2 – Henri Moissan at his bench (left), and his original fluorine cell (right)
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Fluorine: a key enabling element in the nuclear fuel cycle Table I
International fluorspar reserves and production (Roskill, 2009)
China Mexico South Africa Mongolia Russia France Kenya Morocco Namibia Spain USA Other countries Total
2450 630 227 200 200 105 98 95 81 130 0 310 4550
2003 2450 650 240 190 200 110 100 95 85 135 0 320 4540
➤ 1949 sees Simons’ discovery of electrochemical fluorination ➤ In the mid-1950s 3M invents ScotchguardTM ➤ Fried’s initial pioneering work in ‘medicinal fluorine chemistry’ commences in 1954 ➤ Neil Bartlett’s discovery of noble gas chemistry in 1962 (XePtF6) ➤ Rowland and Molina’s model for ozone depletion is published in 1974 ➤ Hargreave’s ‘direct’ perfluorination discoveries in 1979 ➤ Fluorocarbon gases start finding application in the semiconductor industry in late 1980s ➤ 2003 sees O’Hagan’s isolation of first fluorinating enzyme ➤ Fluorine has become ubiquitous in pharmaceuticals, and is essential in medicinal chemistry. At present there are more than 50 industrial producers of HF (Roskill, 2009), the source precursor for all industrial fluorochemicals. Because of cost considerations, it is always preferable in practice to employ a synthesis route that uses HF rather than F2.
The nuclear fuel cycle As indicated in Figure 3, fluorine plays a critical role in several of the unit processes in the nuclear fuel cycle. Generally, the uranium arrives at the conversion plant in the form of U3O8. In order for isotope separation to be effected, uranium is required in the form of a gaseous compound. This compound is UF6. U3O8 is converted to UF6 in a three-step process, each requiring its own plant. The oxide is first converted to UO2 in a hydrogen atmosphere, according to the reaction [3] Note that U3O8 is a mixed valence oxide, thus reduction of a single U+6 to U+4 takes place. Subsequent to this, the uranium dioxide is converted into uranium tetrafluoride in the substitution reaction: [4]
The Journal of The Southern African Institute of Mining and Metallurgy
World reserves (Mt) Reserve
Reserve base
21 41 32 12 10 2 NA 3 6 0 110 230
110 80 40 16 18 14 3 NA 5 8 0 180 480
China South Africa Mexico Mongolia Russia France Kenya Morocco Namibia Spain USA Other Spain
Finally, the uranium tetrafluoride is fluorinated to the hexafluoride, using elemental fluorine gas: [5] Enrichment now takes place, with fissionable U235 separated from U238. This is normally done by centrifuge technology. The enriched uranium is used as solid uranium dioxide. There is more than one way of carrying out the reduction. A standard method is in a hydrogen-fluorine flame reactor. The high temperature is needed to initiate the reaction, given as [6] Solid uranium dioxide powder is pressed into pellets which are housed in Zircalloy tubes. These are bundled into fuel elements (in the case of pressurized-water reactors) ready for use. UF6 is itself a powerful oxidant and fluorinating gas. Materials of construction for plants handling UF6 are thus similar to those for plants that have fluorine as reactant or product. Seals, filters, bearings, etc., are machined from various fluoropolymers, predominantly polytetrafluoroethylene (PTFE). Being fully fluorinated, PTFE is resistant to attack by fluorinating agents (Drobny, 2009; Ebnesajjad, 2013). For comprehensive information about the nuclear fuel cycle, the reader is referred to Barré and Bauquis (2007), Kok (2009), Konings (2012), Tsoulfanidis (1996), Wilson (1996), and Yemelyanov (2011).
South Africa’s fluorochemical capability Highlights in the history of South African fluorochemical technology platform (Naidoo, 2015) are listed below. ➤ Nuclear conversion starts at the Atomic Energy Corporation (AEC) (now the South African Nuclear Energy Corporation, Necsa) in the 1960s ➤ Anhydrous hydrogen fluoride (AHF) and fluorine (F2) are required for uranium hexafluoride (UF6) production. AECI acquires the technology in the 1970s ➤ AECI stops producing AHF in 1984 ➤ Necsa commissions an HF plant in 1985 VOLUME 115
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World mine production (1000 t/a) 2002
Fluorine: a key enabling element in the nuclear fuel cycle
Figure 3 – The nuclear fuel cycle (World Nuclear Association, 2015)
➤ Industrialization and commercialization – 1992 ➤ The Fluorochemical Expansion Initiative (FEI) is adopted as a national initiative – 2006 ➤ Pelchem mandated to champion FEI by Necsa Board of Directors – 2007 ➤ Two research chairs are founded via the National Research Foundation’s SA Research Chairs Initiative (SARChI), one at the University of KwaZulu-Natal (UKZN) and one at the University of Pretoria (UP). At present Necsa and its wholly-owned subsidiary Pelchem SOC Ltd are the main centres of South African fluorochemical expertise, along with two university chairs, one at the University of KwaZulu-Natal (UKZN) and the other at the University of Pretoria (UP). Pelchem runs a commercial 5000 t/a HF plant. Figure 4 shows a photograph of the rotary kiln HF reactor. The company also operates some 20 fluorine electrolysis cells (Figure 5). Pelchem supplies a range of fluorine products to the local and international markets. These include xenon difluoride, nitrogen trifluoride, various organofluorine compounds, perfluorinated alkanes, and a variety of inorganic fluoride salts. Although a full fuel cycle existed on the Pelindaba site, it was abandoned in 1995. The technology in effect does not exist anymore, and if a new fuel cycle is to be established in South Africa, it will have to be a start-up from scratch rather than resuscitation of the old technology. Should this come to pass, our fluorochemical expertise, both existing and under development, will be invaluable if not essential. This will be
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Figure 4 – HF rotary kiln at Pelchem SOC Ltd The Journal of The Southern African Institute of Mining and Metallurgy
Fluorine: a key enabling element in the nuclear fuel cycle
Figure 5 – Pelchem fluorine cells
the case whether the conversion is purchased off the shelf or developed locally. The operation of a conversion plant requires detailed and extensive fluorochemical expertise.
References AGARWAL, A. 2004. Nobel Prize Winners in Chemistry (1901-2002). APH Publishing,, New Delhi, India. AMEDURI, B. 2011. History of fluorine chemistry. Personal communication.
Since the inception of the Fluorochemical Expansion Initiative (FEI), South Africa has made considerable inroads into the development of its fluorochemical capability. The Necsa research effort has been strengthened and expanded, the SARChI Chair at UKZN has been extremely productive regarding research into various thermodynamic aspects of industrial fluorochemical processes, UP is developing a fluoropolymer capability, and Pelchem SOC Ltd has commissioned a new pilot plant, known as the Multipurpose Fluorination Pilot Plant (MFPP). The next few years are critical. A number of things need to happen for the South African research and development effort to continue progressing, and for the technology to be leveraged for the nuclear build plan. These are: ➤ The next phase of FEI needs to be commercially successful, with visible new products ➤ The new cohort of scientists and engineers, trained via funding by FEI, SARChI, and the Advanced Metals Initiative (AMI), have to find employment in the fluorine/nuclear industry ➤ The decision about the nuclear new-build programme has to be taken sooner, rather than later. Within the next 5–7 years the majority of the last generation of Necsa senior scientists and engineers will have retired ➤ The current postgraduate training programme has to be accelerated, with Necsa senior scientists retained for co-supervision of dissertations and theses.
Acknowledgements The author acknowledges the South African National Research Foundation for financial support via the SARChI programme, and Department of Science and Technology for funding through their Fluorochemical Expansion Initiative. Rajen Naidoo, current acting CEO of Pelchem, is thanked the plant photographs, and Dr Johann Nel and Gerard Puts are acknowledged for comment and assistance with graphics, respectively. The Journal of The Southern African Institute of Mining and Metallurgy
BANKS, R.E., SAMRT, B.E., and TATLOW, J.C. (eds). 1994. Organofluorine Chemistry; Principles and Commercial Applications. Springer Science, New York. BARRÉ, B. and BAUQUIS, P.R. 2007. Nuclear Power: Understanding the Future. Ronald Hirlé, Strabourg and Paris. BERTOLINI, J.C. 1992. Hydrofluoric acid: A review of toxicity. Journal of Emergency Medicine, vol. 10, no. 2. pp. 163-168. COTTON, F.A., WILKINSON, G., MURILLO, C.A., and BOCHMANN, M. 2007. Advanced Inorganic Chemistry, Wiley, New Delhi, India. DROBNY, J.G. 2009. The Technology of Fluoropolymers, CRC Press, Boca Raton, FL. EBNESAJJAD, S. 2013. Introduction to Fluoropolymers: Materials, Properties, Applications. Elsevier, Amsterdam. KOK, K.D. 2009. Nuclear Engineering Handbook. CRC Press, Boca Raton, FL. KONINGS, J.M. 2012. Comprehensive Nuclear Materials Vol 2. Elsevier, Amsterdam. NAIDOO, R. 2015. History of fluorine and HF at Necsa. Personal communication. Pelchem SOC Ltd. ROSKILL INFORMATION SERVICES. 2009. The Economics of Fluorspar. London. RUDGE, A.T. 1971. Preparation of elemental fluorine by electrolysis, Introduction to Electrochemical Processes. Kuhn, T. (ed.). Elsevier, Amsterdam. Chapter 5. SHIA, G. 2004. Fluorine. Kirk-Othmer Encyclopedia of Chemical Technology. Hoboken, NJ. SLESSER, C. and SCHRAM, S.B. (eds). 1951. Preparation, Properties, and Technology of Fluorine and Organic Fluoro Compounds. McGraw-Hill, New York. SMITH, R.A. 2004. Hydrogen fluoride. Kirk-Othmer Encyclopedia of Chemical Technology. Hoboken, NJ. TSOULFANIDIS, N. 2010. The Nuclear Fuel Cycle. American Nuclear Society, Scientific Publications, La Grange Park, IL. WILSON, P.D. 1996. The Nuclear Fuel Cycle: From Ore to Waste. Oxford Scientific Publications, Oxford. WORLD NUCLEAR ASSOCIATION. 2015. The nuclear fuel cycle.. http://www.worldnuclear.org/info/Nuclear-Fuel-Cycle/ [Accessed July 2015]. YEMELYANOV, V.S. and YESVSTYUKHEN, A.I. 2011. The Metallurgy of Nuclear Fuel: Properties and Principles of the Technology of Uranium, Thorium and Plutonium. Pergamon Press, Oxford. ◆ VOLUME 115
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a6
Titanium and zirconium metal powder spheroidization by thermal plasma processes by H. Bissett, I.J. van der Walt, J.L. Havenga and J.T. Nel
Synopsis New technologies used to manufacture high-quality components, such as direct laser sintering, require spherical powders 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. Two new metal powder processes have been developed over the past few years. Necsa produces zirconium powders via a plasma process for use in the nuclear industry, and the CSIR produces titanium particles for use in the aerospace industry. Spheroidization and densification of these metal powders require remelting of irregular shaped particles at high temperature and solidifying the resulting droplets by rapid quenching. Spherical metal powders can be obtained by various energy-intensive methods such as atomization of molten metal at high temperatures or rotating electrode methods. Rapid heating and cooling, which prevents contamination of the powder by impurities, is, however, difficult when using these methods for highmelting-point metals. For this reason plasma methods should be considered. Thermal plasmas, characterized 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 suitable for spheroidization of metal powders with relatively high melting points. Thermal plasmas for this purpose can be produced by direct current (DC) plasma arc torches or radio frequency (RF) inductively coupled discharges. In order to obtain chemically pure spheroidized powder, plasma gases such as N2, H2, O2, and CH4 cannot be considered, while Ar, Ne, and He are suitable. Neon is, however, expensive, while helium ionizes easily and it is therefore difficult to obtain a thermal helium plasma at temperatures higher than 3000 K. Therefore argon should be used as plasma gas. Residence times of particles in the plasma region range from 5–20 ms, but this is usually sufficient as 7–8 ms is required for heating and melting of titanium or zirconium metal particles in the 30 μm size range at 3500 K. In this study the melting and spheriodization of titanium powders was investigated by DC non-transferred arc and RF induction plasma methods. The powders were characterized before and after plasma treatment by optical microscopy and scanning electron microscopy (SEM) to observe if any melting or spheroidization had occurred. Keywords plasma, zirconium, titanium, spheroidization.
and Revington, 2010). Direct selective laser sintering (SLS) or, more specifically when considering the processing of metal, direct metal laser sintering (DMLS), are some of these technologies. Implants used in the medical industry to repair or replace bone structure must be strong, ductile, and biocompatible. Materials used for this purpose are stainless steel, cobalt, chromium alloys, or pure titanium, with titanium (Ti) being the material of choice (Taylor, 2004; Bertol et al., 2010). Titanium is also used for applications in the aerospace industry due to its light weight, excellent corrosion resistance, high strength, and attractive fracture behaviour (Li et al., 1997). In order to manufacture a highquality component such as an implant, the titanium powder used as feed material should be dense and spherical. Powder particle size can affect the material spreading and the sintering rate due to the fact that the shape, density, and size of the particles have an effect on their packing density, sintering mechanism, and the flowability of the powder during feeding (Gignard, 1998, p. 34; Despa et al., 2011). The chemical purity of the titanium powder during processing is also very important. Surface oxidation should be prevented as this increases the surface tension, hindering material from flowing during sintering. Oxidation also results in poor bonding between sintered lines affecting the manufactured structures, while nitrides reduce the material’s corrosion resistance (Gignard, 1998, p.34). South Africa has an opportunity to gain benefit from its abundant titanium-bearing mineral reserves through value beneficiation. The potential applications and markets for titanium are in aerospace, the armaments
Introduction
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Contemporary product design processes have been affected by new technologies that assist the manufacturing sector to meet the specifications of specialized components used in the aerospace and medical industries (Williams
Titanium and Zirconium metal powder spheroidization by thermal plasma processes industries, naval applications, offshore oil industries, architecture, and medicine. By establishing a local titanium industry in South Africa, enterprise development across all segments is expected (Van Vuuren, 2009). A convenient way to produce spherical particles, and in particular titanium particles, is to re melt irregularly shaped titanium particles at high temperatures and solidify the droplets by rapid quenching. This can be done by using thermal plasmas, which are characterized by their extremely high temperatures (3 000–10 000 K) and rapid heating and cooling rates (approx. 106 K/s) under oxidizing, reducing, or inert conditions. Plasma is a partially or fully ionized gas in a condition of quasi-neutrality. Thermal plasmas used for particle spheroidization are generally produced by devices such as direct current (DC) arc plasma torches and radio frequency (RF) inductively coupled discharges. Thermal plasmas with temperatures as high as 2 × 104 K can be obtained (Li and Ishigaki, 2001). This paper presents a short background on the DC plasma and RF induction plasma systems used for spheroidization, as well as some background on the properties of powder before and after plasma treatment. A few experimental results relating to the spheroidisation of titanium powder using DC plasmas and RF induction plasmas at the South African Nuclear Energy Corporation (Necsa) are discussed.
Thermal plasmas for spheroidization Although various plasma methods can be used, RF induction plasmas are the preferred method for the spheroidization and densification of particles. This is due to the longer residence time in the plasma and also the lower possibility of contamination caused by electrode erosion. The residence times commonly employed for particle melting in RF plasmas range from 5 to 20 ms (Gignard. 1998, p. 24). DC arc plasma torches, although not regularly used for spheroidization, are also an option. Necsa has expertise in the design and operation of these plasma torches. A short description of the RF and DC plasma torch methods is given, but the complete system, including as particle feeding, quenching, and collection methods, will not be discussed. Information on the spheroidization of titanium metal powders by thermal plasmas has not been found in the public domain, and literature available on spheroidization by DC thermal plasmas is limited to in house designed plasma torches.
Figure 1 – RF induction plasma within the confinement tubes with an induction coil
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RF induction plasma In an induction torch, the energy coupling between the electric generator and the plasma itself is done by a cylindrical coil (Gignard, 1998, p. 6). A typical induction setup is illustrated in Figure 1. The ‘flame’ properties for a RF induction plasma are dependent on the power input, frequency, and the gas pressure and composition. Less power is required at higher frequencies and lower gas pressures to maintain a plasma. Comparing argon and hydrogen, less power is required to maintain an argon plasma due to the fact that argon ionizes more easily than hydrogen (Gignard, 1998, p. 9), while higher plasma gas temperatures can be obtained at higher operating pressures. Although various RF induction plasma systems are available, several research studies have utilized those available from Tekna Systems Incorporated. Most of these studies have used the Tekna PL 50 torch, which can be operated at various frequencies (0.3, 2, or 3 MHz) and power inputs (30, 40, 50 and 100 kW). For spheroidization, Ar or Ar/H2 is used as plasma gas (Jiang and Boulos, 2006; Li and Ishigaki, 2001; Gignard, 1998, pp. 35-42). The use of the Tekna PL 035LS torch has also been reported (Károly and Szépvölgyi, 2005).
DC plasma There are two types of DC torches, these being nontransferred and transferred arc torches. In non-transferred arc DC torches, the electrodes between which the arc is created are inside the body of the torch itself. The plasma gas moves through the arc and is ionized to form the plasma jet. In Figure 2 a typical configuration for a non-transferred arc plasma torch is shown. Non-transferred arc plasma torches are also used in plasma spraying. In this method a powder is introduced into the plasma jet. In this instance, however, spherical powders with a narrow size distribution are used to spray dense, even coatings onto substrates. The purpose of this method is not to spheroidize the powder during the process, but the principle remains similar to spheroidization. In both instances the powder will be quickly melted followed by rapid quenching.
Powder properties Properties such as particle size, bulk density, morphology, impurities levels, etc. are pivotal in selecting a suitable powder to be used in SLS. Particle size affects the material spreading and the sintering rate. The morphology, density, and size of the particles have an effect on their packing
Figure 2 – A typical non-transferred arc plasma torch The Journal of The Southern African Institute of Mining and Metallurgy
Titanium and Zirconium metal powder spheroidization by thermal plasma processes density, sintering mechanism, and the flowability of the powder during feeding (Gignard, 1998, p. 34; Despa and Gheorghe, 2011). Thermal treatment of a powder results in spherical particles with an increased bulk density and decreased in impurity levels, while improving the flowability of the powder.
Table I
Spheroidization of titanium powder in a Tekna argon plasma torch Properties Tap density Production rate Particle size
Density and flowability of powders Bulk density refers to the density of an uncompacted mass of powder taking into account the interparticular voids. The tap density refers to density of a compacted mass of powder where efforts have been made to eliminate the interparticular voids by repeated tapping of the powder. A standardized Hall flow method is used to investigate the apparent density and flow characteristics of free-flowing metal powders (ASTM. 2006). In Table I the increases in tap density and Hall flow rate are indicated for titanium powder treated by a Tekna torch operated at 50 kW (Boulos et al., 2011).
Impurities levels Titanium becomes brittle due to its affinity for oxygen, nitrogen, and hydrogen. Contamination of titanium by air (more specifically oxygen) and hydrogen is thus a problem during welding or sintering. This contamination causes an increase in tensile strength and hardness but reduces ductility, resulting in crack formation. For welding, 0.3% oxygen, 0.15% nitrogen, and 150 ppm hydrogen are seen as the maximum tolerable limits. Surface discolouration gives a good indication of the degree of atmospheric contamination. The colour of the metal changes from silver to a light straw colour (shades of yellow), then dark straw, dark blue, light blue, grey and finally white (TiO2) as contamination increases. The light and dark straw colours indicate light contamination, which is usually acceptable. Dark blue indicates more contamination, while light blue, grey and white indicate high levels of contamination (TWI, 2014). Impurity levels of elements such as O, C, H, N, and S can be determined by using combustion methods or instrumental gas analysis (IGA). In Table I the decrease in impurity levels is clearly indicated for titanium powder treated by a Tekna torch operated at 50 kW using an argon plasma.
Visual characterization of powders Particle morphology can be investigated by using optical microscopy or scanning electron microscopy (SEM) to determine whether melting or spheroidization of the powder
Hall flow test [O] ppm [C] ppm [H] ppm [N] ppm [Cl] ppm [Al] ppm
Before plasma treatment
After plasma treatment
1.0 g/cm3 -140 + 400 mesh / 40 100 μm No flow 2050 160 166 90 1200 80
2.6 g/cm3 10 kg/h at 100 kW -140 + 400 mesh / 40 100 μm 37s/50 g 1800 120 114 90 360 58
occurred during thermal plasma treatment. Figure 3 shows SEM images of titanium powder before and after treatment by a Tekna torch operated at 50 kW using an argon plasma. The images clearly indicate that the rapid heating followed by cooling resulted in spheroidization of the particles.
Experimental The titanium powder used in this study was obtained from the Council for Scientific and Industrial Research (CSIR). Due to the limited quantity of powder available, the as received powder was characterized using SEM and an in-house manufactured Hall flow meter according to ASTM (ASTM. 2014). The amount of powder available for experimentation was not sufficient for post-plasma measurement of the flowability using the Hall flow funnel, and for this reason, SEM and optical microscope images of the powder before and after plasma treatment were used to assess the success of the spheroidization. Similarly, no assessment of the impurity levels of the powder before and after plasma treatment was possible.
As-received powder characterization The as-received and plasma-treated powders were characterized by SEM using a Quanta FEI 200 D instrument. Optical images of the powders were also obtained where possible. A Zeiss Discovery V20 stereo microscope was used for this purpose, and the images recorded utilizing Zeiss Axiovision software. Owing to the large variation in the sizes of the particles observed, the as-received powder was separated into four
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Figure 3 – SEM images of titanium powder (A) before and (B) after plasma treatment
Titanium and Zirconium metal powder spheroidization by thermal plasma processes fractions using sieve mesh sizes of 75, 125, 250, and 425 μm. The four fractions were classified as < 75 μm, 75-125 μm, 125-250 μm and 250-425 μm.
Spheroidization by thermal plasmas Two high-temperature plasma systems were available for the experiments, namely a RF induction plasma and a nontransferred arc DC plasma.
RF thermal plasma system The RF plasma system is shown in Figure 4. This system was used for conceptual work and therefore the exact conditions (gas flow rates, powder flow rates, and input power) of the experiments are not given at this stage. The RF induction coil was connected to a 3 MHz power supply with a maximum power input of 10 kW. Inside the induction coil, two quartz tubes of different diameters were mounted in such a manner as to allow deionized cooling water to flow between them in order to cool of the quartz tube. Above the quartz tube, a hopper loaded with titanium powder was mounted and connected to a 1 mm orifice. Once the system was leak-tight, the system was evacuated and a non-thermal plasmainitiated. Argon gas was slowly introduced. The plasma was gradually changed from a non-thermal to a thermal plasma by increasing the operating pressure (increasing argon flow) and input power until a stable thermal plasma was maintained. At this stage the titanium powder was fed through the thermal plasma. Once a sufficient amount of a given powder fraction had been treated, the plasma was extinguished and the reactor allowed to cool. The powder was then removed for imaging by SEM and optical microscopy. The temperature of the plasma gas was not estimated due to the complexity of the temperature determination method. All four powder fractions were treated.
(resistance) cannot be increased sufficiently. In order to obtain high enough temperatures using argon or helium, the power supply would need to be operated at currents between 400 and 600 A, which fall outside the window of operation of the power supply used in this study. The plasma torch was mounted in a water-cooled reactor and operated at 150 A and 200 V. The plasma gas was fed through the torch at a rate of 1.43 g/s. The plasma gas temperature was estimated by calculating the plasma gas enthalpy (kJ/kg) and relating this to temperature. Once the plasma was stable, the as received powder was fed near the tail flame (plasma jet) of the plasma at a rate of 5 g/min.
Results and discussion As-received powder characterization Figure 5 shows an SEM image of the as-received powder. It is clear that the powder consisted of various particles sizes and morphologies. Some particles appeared to be crystalline while others appeared amorphous. The particles had rough needlelike structures, porous round structures, and porous irregular structures. The SEM images of the four size fractions of the powder are shown in Figure 6. Even after size classification, large variations in particle morphology and structure were observed.
RF thermal plasma treatment
A schematic of the non-transferred arc DC plasma system is shown in Figure 2. The system utilized a 30 kW DC power supply with a maximum current input of 150 A. The system was operated at pressures slightly higher than atmospheric. In order to obtain temperatures high enough for spheroidization of titanium (approx. 3300 K) the system was operated using a nitrogen plasma. The reason for this is that argon or helium ionize easily and therefore the voltage
The temperature of the RF thermal argon plasma could not be calculated or determined at this stage. It is estimated that the temperature of the plasma was near 3300 K (enthalpy of 1.6 MJ/kg for argon) due to the fact that particle melting was achieved. Figure 7 shows an SEM micrograph of the as-received powder after RF thermal plasma treatment. It can be seen that the smaller particles were affected by the thermal treatment. Some spherical particles were observed, as indicated by the arrows. Figure 8 shows a SEM image of the < 75 μm powder fraction after plasma treatment. In most instances it appeared as though spherical particles smaller than 150 μm were obtained, although some agglomeration occurred resulting in large irregularly-shaped particles.
Figure 4 – The RF induction plasma system
Figure 5 – SEM image of the CSIR Ti metal powder
DC thermal plasma system
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Titanium and Zirconium metal powder spheroidization by thermal plasma processes
Figure 6 â&#x20AC;&#x201C; SEM images of the < 75 Îźm, 75-125 Îźm, 125-250 Îźm and 250-425 Îźm powder fractions
DC thermal plasma treatment The DC plasma torch efficiency was calculated to be 61%. Therefore 18.3 kW of the rated 30 kW was available for the plasma gas, which at a nitrogen gas feed rate of 1.43 g/s relates to a gas enthalpy of 12.8 MJ/kg for nitrogen. From thermodynamic data it is estimated that the gas temperature near the arc was approximately 6200 K. Further from the arc the temperature was approximately 3300 K (enthalpy of 3.7 MJ/kg for nitrogen). For the DC plasma treatment, only the as-received powder was used and not the separate fractions. The estimated enthalpy cost per unit mass titanium for spheroidization was 60 kWh/kg (216 MJ/kg). Optical micrographs of the treated powder are shown in Figure 9. A high degree of spheroidization is evident. Due to the fact that nitrogen was used as plasma gas, surface contamination occurred, which is evident from the colours
Figure 7 â&#x20AC;&#x201C; SEM image of the as-received powder after RF thermal plasma treatment. The arrows indicate spheroidized particles
Figure 8 â&#x20AC;&#x201C; SEM image of the < 75 Îźm fraction after RF thermal plasma treatment
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Similar results were obtained using the 75 125 Îźm fraction, although slightly fewer spherical particles were observed. The large fractions (125-250 Îźm and 250-425 Îźm) were not affected by the plasma treatment. These results suggest that particles smaller than 125 Îźm can be spheroidized. Agglomeration of melted particles is a concern, however, and for this reason a sheath gas is required. Agglomerate formation is possibly due to the collection of fine particles near the quartz tube. The particles â&#x20AC;&#x2DC;rollâ&#x20AC;&#x2122; down the tube surface, forming relatively large agglomerates. The sheath gas is not only necessary to prevent agglomeration from occurring, but will also prevent the collection of fine particles near the coil region of the quartz tube, a regularly observed occurrence. Once titanium metal particles collect on the surface of the quartz tube, induction heating of the particles occurs, resulting in a reduction of the plasma gas temperature.
Titanium and Zirconium metal powder spheroidization by thermal plasma processes
Figure 9 – Optical micrographs of the DC thermal plasma treated powder
observed in the images. The light and dark straw colours indicate light contamination, which is usually acceptable. Dark blue indicates more contamination, while light blue indicate high levels of contamination (TWI, 2014). A SEM image of the treated powder is shown in Figure 10. The variation in the contrast of the particles indicates various degrees of contamination by either nitrogen or oxygen. Again, a high degree of spheroidization is evident. A nitrogen plasma was used in this instance, and it is expected that an argon plasma will yield similar results with respect to spheroidization of the particles, but without surface contamination occurring.
Conclusions Available literature indicated that titanium metal particles can be spheroidized in a thermal plasma by re melting of irregularly shaped particles at high temperatures and solidifying the droplets by a rapid quench. Thermal plasma treatment of powders results in improved powder flow characteristics and an increased density of individual particles, pivotal in selecting a suitable powder to be used in SLS or DMLS. The titanium powder received from the CSIR consisted of various particle sizes and morphologies, where some particles appeared to be crystalline and others amorphous. In this study SEM images indicated that the powder could be spheroidized by an RF thermal plasma using argon gas without increasing impurity levels significantly. Various powder fractions were tested. Only the < 75 μm and 75-125 μm fractions could be spheroidized; the larger fractions (125250 μm and 250-425 μm) were not affected by the plasma treatment. SEM and optical microscope images also showed that titanium powder could be effectively spheroidized by a DC thermal plasma using nitrogen gas. In this instance the powder was used as-received and not sieved into fractions. The particles changed colour, indicating surface contamination by nitrogen. It is, however, expected that spheroidization without contamination will be possible by DC thermal plasma treatment using argon gas at a gas temperature near 6200 K.
References ASTM INTERNATIONAL. 2006. Standard test method for apparent density of freeflowing metal powders using the Hall flowmeter funnel. Designation: B212 99. West Conshohocken, PA. BERTOL, L.S., JÚNIOR, W.K., DA SILVA, F.P., and AUMUND, K.C. 2010. Technical report Medical design: Direct metal laser sintering of Ti 6Al 4V. Materials and Design, vol. 31. pp. 3982 – 3988.
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Figure 10 – A SEM image of the DC thermal plasma treated powder
BOULOS, M., HEBERLEIN, J., and FAUCHAIS, P. 2011. Thermal plasma processes, fundamentals and applications. Short course presented at the University of Pretoria, 28 29 October. 2011. Pretoria, South Africa. 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. GIGNARD, N.M. 1998. Experimental optimization of the spheroidization of metallic and ceramic powders with induction plasma. Thesis, National Library of Canada, Sherbrooke, Quebec, Canada. . JIANG, X. and BOULOS, M. 2006. Induction plasma spheroidization of tungsten and molybdenum powder. Transactions of Nonferrous Metal Society of China, vol. 16. pp. 13 – 17. LI, Z. GOBBI, S.L. NORRIS, I. ZOLOTVSKY, S., and RICHTER, K.H. 1997. Laser welding techniques for titanium alloy sheet. Journal of Materials Processing Technology, vol. 65. pp. 203 – 208. 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. TAYLOR, C.M. 2004. Direct laser sintering of stainless steels: thermal experiments and numerical modelling. PhD thesis, School of Mechanical Engineering, University of Leeds, UK. Chapter 2, p. 5. TWI. http://www.twi-global.com/technical-knowledge/job-knowledge/weldingof-titanium-and-its-alloys-part-1-109 [Accessed 18 July 2014]. VAN VUUREN, D.S. 2009. Titanium – an opportunity and challenge for South Africa. Keynote address: 7th International Heavy Minerals Conference ‘What Next’, Champagne Sports Resort, Drakensberg, South Africa, 20–23 September 2009. Southern African Institute of Mining and Metallurgy, Johannesburg. pp. 1 – 8. 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 Association of Oral and Maxillofacial Surgeons, vol. 39. pp.182 – 184. ◆ The Journal of The Southern African Institute of Mining and Metallurgy
http://dx.doi.org/10.17159/2411-9717/2015/v115n10a7
Plasma technology for the manufacturing of nuclear materials at Necsa by I.J. van der Walt, J.T. Nel and J.L. Havenga
The development of plasma technology at Necsa started in the early 1980s, when the applicability of high-temperature plasmas in the nuclear fuel cycle was investigated. Since 1995, this plasma expertise has expanded to other industrial applications, for example mineral beneficiation, nanotechnology, fluorocarbon production and waste treatment, all of which are also of relevance to the nuclear industry. Necsa has demonstrated the manufacture of plasma-dissociated zircon, zirconium metal powder, carbon nanotubes, silicon carbide (SiC), zirconium carbide (ZrC) and boron carbide (B4C) at the laboratory and pilot plant scale. These materials are commonly used in the nuclear industry. Zirconium alloys are used as fuel cladding material for nuclear fuel assemblies. Necsa manufactured the monomer tetrafluoroethylene (TFE), using 150 and 450 kW DC plasma systems, from which the polymer polytetrafluoroethylene (PTFE) was synthesized for use in filters and as seals in nuclear plants. With the nuclear renaissance at hand, it was demonstrated that plasma technology can be used to produce hydrofluoric acid (HF), which is used in the manufacture of fluorine gas (F2) for the production of uranium hexafluoride (UF6) directly from the mineral calcium fluoride (CaF2) without the use of sulphuric acid as in the conventional process. The recovery of valuable uranium from nuclear waste such as filters, oils, and solids with plasma processes will also be discussed. The destruction of low-level nuclear waste by a plasma gasification system can reduce the volume of this waste by several orders in magnitude, resulting in huge savings in the storage costs. Another product of plasma technology is the encapsulation process for nuclear waste and the production of vitrified product, which could be used as filler material for medium-level nuclear waste. Keywords plasma, zirconium, fluorocarbons, nuclear waste, nanomaterials.
Introduction During the 1980s, a plasma research and development programme was launched at the erstwhile Uranium Enrichment Corporation of South Africa (UCOR) at Valindaba outside Pretoria. The main purpose of this programme was to develop alternative and more economical processes for the processing and manufacturing of uranium compounds used in the nuclear fuel cycle. The conversion of UO2 (as received from uranium mines) to UF6, which is needed for enrichment, conventionally involves several laborious, expensive, chemical processing steps. Plasma processes can eliminate several of the intermediate steps. The Journal of The Southern African Institute of Mining and Metallurgy
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Synopsis
Enriched UF6 is converted to UO2, which is pressed into pellets that are used in nuclear power plants. Plasma technology could also be used in a modified and optimized conversion process. Zirconium alloys are used as nuclear fuel cladding material. In the 1980s, UCOR developed plasma process to produce zirconium metal from ZrCl4, using hydrogen as a reductant (Nel et al., 2011). With the termination of South Africa’s nuclear programme in the 1990s, the South African government asked the then Atomic Energy Corporation (AEC) to investigate the possible re-aligning of already established plasma technology, expertise and equipment to investigate other applications for plasma technology in the non-nuclear industry. From this the so-called Metox (Metal Oxide) programme was born. This programme investigated alternative means of zircon beneficiation, as well as the production of high-temperature-resistant ceramic compounds that are used in the nuclear industry, such as alumina, zirconia, silica, silicon carbide, zirconium carbide, and boron carbide. By changing the process parameters, nanoparticles of the abovementioned products were also synthesized in a plasma system. Polytetrafluoroethylene (PTFE) was extensively used in nuclear plants as seals, in valves and pipes, as containers, filters, etc. PTFE filters used in the enrichment plant that are contaminated with uranium products can be destroyed with a plasma process and simultaneously, in the same process, the uranium values can be recovered. There is no PTFE or fluoropolymer production facility in South Africa and all these products are imported. In 1994, Necsa started
Plasma technology for the manufacturing of nuclear materials at Necsa a project on the production of TFE (C2F4), the precursor for PTFE, using a plasma process. South Africa is the world’s fourth-largest producer of calcium fluoride (fluorspar or CaF2), but very little or no local beneficiation takes place. In the nuclear fuel cycle, hydrofluoric acid is produced from CaF2 to manufacture fluorine gas, which is used to make UF6 for the enrichment purposes. It was proved that HF can be manufactured directly in a plasma system instead of using the conventional sulphuric acid route. The destruction, encapsulation, and vitrification of nuclear waste can also be accomplished with a plasma process. Under the Advanced Metals Initiative (AMI) programme of the Department of Science and Technology (DST), a new process was developed to make nuclear-grade zirconium metal in a continuous plasma process. The purpose of this paper is to discuss the abovementioned processes in more detail.
Plasma technology in the nuclear fuel cycle The nuclear fuel cycle is schematically presented in Figure 1. The red arrows indicate where plasma technology can be used to replace conventional processes. UF6 is a gas that is used for the enrichment of uranium. Processed uranium ore as received from the mines may consist of various chemical compounds, depending on the specific process that the mine uses. The conventional conversion route for UF6 production is a complex process involving multiple steps, each of which requires a separate, complete chemical plant. There are many minor variations on this process, but in general it can be described in as follows. Uranium oxide, which can be UO3 or U3O8, (also referred to as yellowcake) or ammonium diuranate (ADU or (NH4)2U2O7) is calcined to UO2 with hydrogen. The UO2 is
hydrofluorinated with hydrofluoric acid (HF(g)) to form UF4, a solid green compound, which is then fluorinated with fluorine gas to form UF6(g). This process is schematically presented in Figure 2. Hydrofluoric acid is produced by the reaction of CaF2 with H2SO4 according to Equation [1]. F2 gas is produced by electrolysis of KF.HF. CaF2 + H2SO4 → CaSO4 + 2HF(g)
[1]
The AEC proved in concept that a uranium oxide could be directly fluorinated to UF6 with capacitive and inductive coupled non-thermal plasma according to Equation [2] (Jones, Barcza, and Curr, 1993). This will eliminate the multistep process presented in Figure 2 with a major economic advantage. 2UO3 + 3CF4 → 2UF6 + 3CO2
[2]
Uranium metal can also be fluorinated by F2 or CF4 or mixtures thereof in an inductive coupled non-thermal plasma (Equation [3]). U(m) + F2/CF4 → UF6
[3]
The AEC (now Necsa), demonstrated the use of a direct current (DC) plasma process to convert UF6 to UF4 using hydrogen or cracked ammonia as reductant. This process was scaled to a production rate of 1 kg UF4 per hour in a continuous operation. It was estimated that the process could be 30% more economical than the conventional process. The manufacture of depleted uranium metal from UF4 was also demonstrated, and a total of 13 kg of depleted uranium was manufactured using this process by the AEC using a scaledup laboratory system. Processes have been proposed by Toumanov (2003) and Fridman (2008, p. 449) for the direct conversion of UF6 to uranium metal (Equation [4]). UF6(gas) → Usolid + 3F2(g)
[4]
Figure 1 – A schematic of the nuclear fuel cycle. (World Nuclear University, n.d.) Red arrows indicate where plasma technology can be used
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Plasma technology for the manufacturing of nuclear materials at Necsa
Figure 2 – The conversion of uranium oxide to UF6
This process is highly endothermic and takes place at temperatures above 5000 K. High quenching rates in the order of 10-7 seconds are required to prevent the reverse reaction from taking place. Beta-UF5(s) is an unwanted reaction product that forms under certain conditions during the enrichment process. A high-frequency non-thermal plasma process was utilized for the in situ back-fluorination of beta-UF5 to UF6 using a mixture of CF4/argon or fluorine/argon. Unfortunately, further developments regarding the abovementioned nuclear plasma processes were halted when South Africa’s nuclear programme was terminated in 1993.
Nuclear nano-materials
4BCl3 + CH4 + 4H2 → B4C + 12HCl The Journal of The Southern African Institute of Mining and Metallurgy
[5]
Figure 3 – Schematic of a typical plasma system for the manufacturing of nano-powders
Nano-sized ZrC can also be manufactured by a plasma process.
Fluorocarbons in the nuclear industry PTFE is one of the few fluorocarbons used in the nuclear industry. A sintered filter was produced from PTFE and used in the enrichment stages. A thermal chemical process was developed to produce CF4 gas by reacting solid carbon with pure fluorine. A 99 % conversion was achieved, and the gas was used in a plasma process to produce PTFE monomer. The chemical formulae governing these reactions are presented in Equations [6] and [7]: C(s) + 2F2(g) → CF4
[6]
CF4 + C(s) + Plasma → C2F4
[7]
Various CF4 plasma plants for the production of TFE (C2F4) were constructed and commissioned at Necsa. A small 30 kW laboratory system and 100, 150, and 450 kW pilot systems were successfully developed and operated. A typical plasma TFE production system is presented in Figure 4. As part of a DST-funded Fluorochemical Expansion Initiative, the C2F4 (TFE) was used in a suspension polymerization process to produce PTFE. After PTFE had been produced successfully, the second stage of the project to produce FEP was developed and proven. A schematic diagram of the system is presented in Figure 5. Since neither of these polymers are produced in South Africa, this an opportunity for a new business. VOLUME 115
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Plasma technology is an extremely useful and effective technique for the manufacture of nanoparticles, especially with regard to advanced ceramics (Fridman, 2008, pp 417–498, 566; Van der Walt, 2012; Boulos, Fauchais and Pfender, 1994). A typical plasma system for the manufacturing of nanopowders is presented in Figure 3. The process involves the evaporation of reactants in the high-temperature environment of a plasma arc, followed by rapid quenching of the vapour in order to nucleate the particles. Ceramic powders such as carbides, nitrides and oxides have been synthesized in this way. Ceramic materials like SiC, ZrC, B4C and ZrO2, which are stable at high temperatures, have numerous applications in the nuclear industry, and will become even more prominent in Generation IV high-temperature gascooled nuclear reactors (HTGRs) (Konings et al., 2012; Kok, 2009). Necsa has produced several of these nanomaterials. Nano-zirconia can be produced by the plasma dissociation of zircon, followed by selective removal of the formed amorphous silica by fluorinating agents. Necsa commissioned and operated a pilot plant with a capacity of 100 kg/h to produce plasma-dissociated zircon and a 10 kg/h plant for the production of nano-silica. Boron carbide is often applied in nuclear technology as a neutron shielding material and as control rods inside nuclear reactors. It has a high neutron absorbance cross-section for thermal neutrons of 755 barns and a high melting point of 2723 K (Kirk-Othmer Encyclopedia of Chemical Technology; Lipp, 1965). Necsa produced nuclear-grade B4C powder in a 30 kW non-transferred arc DC plasma system according to Equation [5]. BCl3 was evaporated at 60°C and fed into the plasma reactor at a rate of 120 litre per hour. The reaction took place at about 2200°C, and nitrogen was used as a quench gas. A particle size of between 80 nm and 100 nm was obtained, depending on the plasma parameters. Nanosized B4C powder has enhanced compaction properties for the production of B4C pellets.
Plasma technology for the manufacturing of nuclear materials at Necsa
Figure 4 – Schematic of a plasma TFE production process
Plasma treatment of nuclear waste Vitrification Vitrification is a recognized standard method of reducing the environmental impact of harmful waste. In vitrification, the solid component of the treated waste is encapsulated by hightemperature treatment. In recent years a glass formulation has been added to the vitrification process that allows the tailoring of the waste matrix. Nuclear waste containment in glass has several advantages: it guarantees the stability of the waste package over very long time periods (50–300 years) and could reduce the waste volume. Disposal safety criteria can be more easily met by adaptation of the waste
matrix. This aspect is crucial in the treatment and disposal of nuclear waste due to the long periods of storage required. With the addition of a glass/ceramic mixture that encapsulates the waste, the resistance to leaching and chemical attack can be optimized (Petijean et al., 2002). The matrix developed for the containment of radionuclides should take into account the following issues (Luckscheiter and Nesovic, 1996; Bisset and van der Walt, 2009): ➤ The waste must be accommodated within the matrix composition ➤ Waste loading ➤ Characteristics such as resistance to solubility, phase separation, and devitrification ➤ Long-term behaviour such as thermal stability, irradiation resistance and chemical and mechanical stability ➤ Process considerations such as melting temperature, viscosity and electrical conductivity. A thermal plasma is able to decompose various types of waste into a gas and a residue by exposing it to a very high temperature. Because of the very high temperature in the plasma, no sorting is necessary, thus human exposure is minimized. This system also has the capability of containing the uranium contamination in the molten pool of metal (if the waste is in a drum) as well as in ceramics and noncombustible materials. The cost implications will include a capital installation cost and an operational cost. If designed properly, a plasma volume reduction system for nuclear waste can generate electricity to feed back into the process, reducing the operating cost to a minimum. There is a case to be made whether the long-term cost of disposal at a waste repository like Vaalputs will be comparable to the costs of volume reduction of nuclear waste by a plasma process. The additional benefits of the plasma system are the production of a vitrified product and the reduction of the waste volume (van der Walt and Rampersadh, 2011).
Figure 5 – Schematic presentation of the polymerization reactor
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Plasma technology for the manufacturing of nuclear materials at Necsa Uranium recovery The recovery of valuable constituents such as uranium from historical and current waste streams is important and may also be economically viable. Plasma technology could replace other competing processes for this purpose, and may have important advantages over the conventional processes. Conceptual research was done on one such waste stream, namely the coated uranium kernels used in the pebble bed modular reactor (PBMR) project. An alternative recovery method to the mechanical crushing of off-specification tri-structural-isotropic (TRISO) coated fuel microspheres was demonstrated. It was shown that the inert SiC layer can be completely removed by etching with the active fluorine species by an inductively coupled radio-frequency CF4 glow-discharge impinging a static bed from the top, at a working pressure of 1 kPa (van der Walt et al., 2011). The apparatus is shown in Figure 6. The recovery of uranium from contaminated PTFE filters was also investigated by Necsa. A process was developed that depolymerizes PTFE and thereby separates the matrix from the extremely valuable uranium. This system makes use of radio frequency (RF) induction heating to heat a reactor up to the depolymerization temperature of PTFE. Inside the reactor the PTFE is depolymerized and the product gas is cooled, separated from residual uranium particles, scrubbed and evacuated to a destruction facility where fluorocarbons are converted into CO2 and HF by means of a DC thermal plasma
system using O2 and water gas. The HF is scrubbed by means of a KOH scrubber. A schematic diagram of this system is presented in Figure 7.
The AMI zirconium process In 2005, the DST initiated the Advanced Metals Initiative (AMI) programme and asked Necsa to coordinate the New Metals Development Network, with the specific focus on the manufacturing of nuclear-grade zirconium metal from the mineral zircon (Zr(Hf)SiO4). The request from DST was specifically to develop a new and more economical way to manufacture Zr metal using Necsa’s plasma and fluorochemical expertise and existing facilities. The conventional methods are described in detail in the scientific and patent literature dating back to the early 1920s (Blumenthal, 1958). Significant industrial manufacturing of nuclear-grade zirconium started in the late 1940s after World War II, when the advantages of using zirconium in nuclear power plants were realized. In general, most industrial manufacturing processes for nuclear-grade zirconium use high-temperature carbochlorination of zircon, selective separation of ZrCl4 and SiCl4, the separation of Zr and Hf by means of solvent extraction, distillation, selective crystallization, etc., and eventually the reduction of purified ZrCl4 with magnesium via the Kroll process (Equation [2]). These conventional processes can consist of between 16 and 18 individual steps, making the manufacturing of nuclear-grade zirconium one of the most expensive operations in the nuclear fuel cycle. ZrCl4 + 2Mg → Zr + 2MgCl2
[8]
Zircon is an extremely chemical inert mineral. However, activating zircon with a plasma process to produce plasmadissociated zircon (PDZ) makes it very reactive, especially towards fluoride-containing compounds such as HF and ammonium bifluoride (ABF, NH4F.HF) (Nel et al., 2011a, 2011b). Necsa operated a semi-commercial PDZ plant from 1995 to 2003, which had a capacity of 100 kg/h using a 3 × 150 kW DC plasma torch configuration (Havenga and Nel, 2011). Necsa developed a plasma process for the manufacturing of Zr metal powder from PDZ, making use of the reactivity of PDZ with ABF to produce ZrF4 (Makhofane et al., 2011). ZrF4 or ZrCl4 is reduced with magnesium in a in a 30 kW DC non-transferred arc plasma plasma process
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Figure 7 – Uranium recovery system with plasma waste destruction VOLUME 115
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Figure 6 – Non-thermal CF4 plasma system
Plasma technology for the manufacturing of nuclear materials at Necsa
Figure 8 â&#x20AC;&#x201C;Block flow diagram of the AMI zirconium metal process
(similar to the Kroll process) (Nel et al., 2012). The whole process is generally referred to as the AMI zirconium metal process (Nel et al., 2013). The plasma process has, however, the additional advantage that it can be modified to make it a continuous process, unlike the conventional Kroll process. The AMI zirconium metal process making use of plasma/fluoride technology consists of only six steps, and offers the potential of huge cost savings in comparison with the conventional processes. This process is schematically presented in Figure 8.
Conclusions Plasma processing has numerous potential applications in nuclear science and technology and for the manufacturing of nuclear materials. Over a period of three decades, Necsa has
developed plasma applications in the nuclear fuel cycle, including the manufacture of nuclear ceramics and nanoparticles that are being (or can be) used in nuclear reactors, especially in high-temperature gas-cooled reactors. Fluoromonomers, which are used as precursors for many fluoropolymers, can also been made via plasma processes. Fluoropolymers are extensively used as filters, seals, and containers in the nuclear industry. Nuclear waste destruction can also be accomplished by plasma processes. Many of these processes have been developed to pilot scale, while others were developed only to the laboratory scale and proof-ofconcept. Necsa has patented plasma and fluoride processes for the manufacturing of nuclear-grade zirconium metal powder using the mineral zircon as starting material. â&#x2014;&#x2020;
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a8
Synthesis and deposition of silicon carbide nanopowders in a microwaveinduced plasma operating at low to atmospheric pressures by J.H. van Laar*†,I.J. van der Walt*, H. Bissett*, G.J. Puts† and P.L. Crouse†
Silicon carbide nanopowders were produced using a microwave-induced plasma process operating at 15 kPa absolute and at atmospheric pressure. Methyltrichlorosilane (MTS) served as precursor, due to its advantageous stoichiometric silicon-to-carbon ratio of unity, allowing it to act as both carbon and silicon source. Argon served as carrier gas, and an additional hydrogen feed helped ensure a fully reducing reaction environment. The parameters under investigation were the H2:MTS molar ratio and the total enthalpy. The particle size distribution ranged from 20 nm upwards, as determined by SEM and TEM micrographs. It was found that an increase in enthalpy and a higher H2:MTS ratio resulted in smaller SiC particle sizes. The adhesion of particles was a common occurrence during the process, resulting in larger agglomerate sizes. SiC layers were deposited at 15 kPa with thicknesses ranging from 5.8 to 15 μm. Keywords silicon carbide, microwave plasma, methyltrichlorosilane, nanoparticles.
Introduction It is well known that methyltrichlorosilane (CH3SiCl3 or MTS) decomposes to form silicon carbide (SiC) and hydrogen chloride (HCl) as shown in Equation [1]. The reaction kinetics and mechanisms for this reaction are thoroughly reported in the literature (Sone et al., 2000; Papasouliotis and Sotirchos, 1999; Kaneko et al., 2002; Wang et al., 2011). CH3SiCl3(l) →SiC(s) + 3HCl(g)
[1]
Owing to its physical and mechanical properties, SiC has found application in several areas of high-power, high-frequency, and high-temperature technology (Harris, 1995; Saddow and Agarwal, 2004). Amongst its seemingly limitless applications, SiC is gaining increased attention as a nuclear ceramic due to its excellent mechanical properties and dimensional stability under irradiation (Katoh et al., 2012). SiC composites have also been proposed to act as fuel cladding in light water reactors as a direct replacement for zirconium alloy cladding (Katoh et al., 2012). Conventional ceramics exhibit certain drawbacks such as low ductility and high brittleness; however, the fact that nanopowders can overcome these shortcomings has encouraged the use of SiC in various fields. The application of SiC The Journal of The Southern African Institute of Mining and Metallurgy
* Applied Chemistry Division, South African Nuclear Energy Corporation (NECSA) SOC Ltd., Pelindaba, South Africa. † Fluoro-Materials Group, University of Pretoria, South Africa. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. VOLUME 115
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Synopsis
nanopowders in the nuclear industry, for example, allows for fast recovery of irradiation-induced defects (Vaßen and Stöver, 2001). Various synthesis methods for SiC nanoparticles have been reported in the literature. These include carbothermic reduction (Dhage et al., 2009), pulsed laser deposition (Kamlag et al., 2001), sol-gel processes (Ahmed and El-Sheikh, 2009), microwave heating (Satapathy et al., 2005; Moshtaghioun et al., 2012) as well as numerous plasma techniques ranging from inductive radio-frequency (RF) (Károly et al., 2011; Sachdev and Scheid, 2001) to microwave plasma-assisted chemical vapour deposition (Tang et al., 2008; Honda et al., 2003; Vennekamp et al., 2011). Previous work by the authors (Van Laar et al., 2015) reported the empirical study of the synthesis of SiC powders using a microwave-induced plasma at atmospheric pressure only. In this work a microwave plasma system similar to that reported in the earlier paper (Van Laar et al., 2015) was used. The study is expanded here to include initial results of SiC deposition at low pressure (15 kPa) onto a quartz substrate. MTS was chosen as precursor due to its advantageous stoichiometric siliconto-carbon ratio of unity, allowing it to act as both a carbon and a silicon source. A combination of its liquid state, high vapour pressure at standard conditions, and high volatility also allows for easier feeding of the MTS into the system at both low and high pressures. Hydrogen was fed into the reactor to serve as a reductant for driving the
Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma conversion reaction. Argon served as the carrier gas. Changes in total enthalpy and the H2:MTS molar ratio were studied at atmospheric pressure.
Experimental Apparatus The experimental set-up consisted of a 1.5 kW power supply with a MOS-FET amplifier, a microwave generator operating at 2.45 GHz, a water-cooled magnetron head, and rectangular waveguide with sliding short and stub tuners. The quartz tube, with internal diameter of 2 cm and a length of 30 cm, was positioned through the middle and perpendicular to the
Figure 1 – Physical layout of the reactor assembly
waveguide and served as the reaction zone. Support flanges at the top and bottom of the tube also served as gas inlets. Argon and hydrogen flow rates were controlled using calibrated Aalborg rotameters. This set-up is similar to the one used previously (Van Laar et al., 2015). An illustration of the physical layout of the reactor assembly is shown in Figure 1, and a schematic representation of the flow path is shown in Figure 2. At atmospheric pressures, argon was bubbled through the MTS in order to help vapourize and carry the MTS through the system. The MTS-rich argon stream was then mixed with hydrogen and argon streams in predetermined ratios before entering the reactor. At low pressures, the MTS was vapourized under vacuum, and the flow rate controlled using a valve placed before the entrance to the reactor. Calibration curves were reported in previous work (Van Laar et al., 2015). The exiting gas was passed through a CaCO3 scrubber to remove HCl and any unreacted MTS before entering the extraction system. A T-connection and valve assembly immediately following the scrubber allowed for shifting between vacuum and atmospheric operating pressures. Characterization of the particles was performed using the available equipment at the University of Pretoria. Particle size distribution was determined with a ZEN 3600 Malvern Zetasizer Nano System. Scanning electron microscopy (SEM) was performed on the particles using a high-resolution (6 Å) JEOL 6000 system, and transmission electron microscopy (TEM) was performed using the JEOL JEM2100F TEM (JEOL Japan). Powder X-ray diffraction was conducted with a PANalyticalX`pert Pro diffractometer using Co Kα radiation. The peaks were assigned using the databases supplied by the instrument manufacturer.
Figure 2 – Schematic diagram of the experimental set-up
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Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma
Figure 3 – Operation of the reactor at (a) atmospheric pressure and (b) low pressure
At the start of each experimental run the argon plasma was initiated under vacuum at approximately 15 kPa (Figure 3b), using an Alcatel 2010I duel-stage rotary vane pump. The deposition experiments were performed at this pressure. Atmospheric pressure was reached by gradually increasing the operating pressure, at which point filamentation of the plasma structure occurred (Cardoso et al., 2009) as shown in Figure 3a. Hydrogen and MTS were then fed into the system. Depending on the H2:MTS ratio, stable plasmas were possible at applied powers between 200 and 1500 W. High hydrogen concentrations tended to extinguish the plasma, presumably due to increased energy demand for dissociation of the H2 bonds. During experiments at atmospheric pressure, a black powder deposited on the inner walls of the quartz tubes. These tubes were removed after each experimental run and flushed using distilled water. The water was collected and evaporated in a drying oven at 80°C, after which the black powders were collected. Quartz tubes with smaller diameters (15 mm) were used as substrates for the deposition experiments. These smaller tubes were placed inside the larger tubes, and held in place using zirconium wool.
Results and discussion Synthesis of SiC at atmospheric pressure The results of the synthesis experiments are reported elsewhere (Van Laar et al., 2015). The enthalpy values are those of the system enthalpy, HT, which combines all the chemical species (argon, hydrogen, and MTS). These enthalpy values were determined by Equation [2]: [2] where Pf is the forwarded power, Pr is the reflected power, • and mT is the total mass flow rate. The enthalpy of the MTS, HMTS, was calculated from Equation [3]: [3]
The Journal of The Southern African Institute of Mining and Metallurgy
The average particle sizes, as determined from SEM and Zetasizer results, are also listed elsewhere (Van Laar et al., 2015). Particle agglomerates were a common occurrence. Based on the results, the best-fitted model included quadratic and 2-factor interaction terms. Analysis of variance (ANOVA) results for the agglomerate sizes indicated that enthalpy had the greatest effect on the agglomerate sizes, whereas the H2:MTS ratio was found to be least significant (Van Laar et al., 2015). Zetasizer and SEM results were analysed using response surface analysis (RSA). The resulting surface contour plots are shown in Figure 4 and Figure 5 respectively. Figure 4 suggests that particle size decreases with increasing enthalpy and H2:MTS ratio. Considering the effect of the two studied parameters separately, higher enthalpy values presumably allow de-agglomeration to occur more readily due to more energetic particle collisions. This results in smaller particle sizes. Higher H2:MTS ratios result in a more reducing environment, leading to smaller particle sizes. When considering the effect of both parameters in conjunction, two contrasting trends are seen. At high enthalpy values (195–220 MJ/kg), particle sizes seem to increase with increasing H2:MTS ratios. This trend could possibly be attributed to the increasing energy demand for hydrogen dissociation with increasing H2:MTS ratio. At low enthalpy values (70 –120 MJ/kg), particle size decreases with increasing H2:MTS ratio. This is in contrast to trends seen at high enthalpy values. It is speculated that this trend at low enthalpy values occurs because the energy supply is not adequate to allow for hydrogen dissociation, increasing the reducing environment and allowing more available energy for de-agglomeration. The particle size distributions determined by the Zetasizer show that lower enthalpies produce larger agglomerates, but that the agglomeration process is much more sensitive to H2:MTS ratios, with higher ratios negatively influencing the agglomerate size. Figure 6 and Figure 7 present SEM micrographs showing agglomerate particle sizes down to approximately 50 nm. VOLUME 115
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Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma
Figure 4 – Effect of total enthalpy and H2:MTS molar ratio on individual particle size
Figure 5 – Effect of total enthalpy and H2:MTS molar ratio on agglomerate size
Figure 6 – SEM image of SiC (1)
Nanoparticles tend to form strongly bound agglomerates. Ultrasound sonification is often used to de-agglomerate nanoparticles; however, there is insufficient knowledge on the de-agglomeration of nanoparticulate systems (Sauter et
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Figure 7 – SEM image of SiC (2)
al., 2008), making it extremely difficult to obtain a reliable particle size distribution. From the images in Figure 6 and Figure 7, it can be seen that agglomerates exist in a wide variety of sizes. The smallest particle sizes to be confidently identified from SEM images were approximately 50 nm. The TEM micrographs are presented in Figure 8 and Figure 9. Figure 8 shows particle sizes down to approximately 20 nm, and Figure 9 shows larger structures, presumably those of the agglomerates. X-ray diffraction results show diffraction peaks at positions indicative of beta (β) SiC (also referred to as cubic) shown in Figure 10. Also present are peaks indicative of silicon, although in much lower amounts than SiC. This could imply the presence of silicon in the SEM and TEM images. Other elements present within the plasma were verified using an optical emission spectrometer, reported previously (Van Laar et al., 2015). The majority of peaks were in good agreement with experimental values of elemental silicon, carbon, and argon (Kramida et al., 2014), suggestive of MTS decomposition in the plasma. The presence of elemental silicon in the gas phase as well as silicon in the product material suggests that the addition of hydrogen to the plasma drives the conversion reactions too far into the reductive regime. The equilibrium thermodynamics and formation mechanisms of Equation [1] have been reported in the literature (Deng et al., 2009). The thermodynamics software package TERRA (Trusov, 2006) was used to confirm the optimum conditions for the formation of β-SiC. The results predict that optimum yield for β-SiC formation is achieved at temperatures of around 1400 K. Microwave plasmas are known to achieve temperatures in the region of 1000–10 000 K (Tendero et al., 2006). The temperatures of the experiments reported in this paper were measured using a pyrometer, giving an indication of the SiC temperature inside the reactor. SiC is chemically inert, with excellent microwave absorption and heat-conducting properties (Isfort et al., 2011). This enabled a rough estimation of the temperatures inside the reactor, which were measured to range from 1100 to 1400 K. The Journal of The Southern African Institute of Mining and Metallurgy
Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma
Figure 8 – TEM image of silicone carbide (1)
Figure 9 – TEM image of silicone carbide (2)
Figure 10 – XRD spectrum of a product sample synthesized at an H2:MTS ratio of 4. SiC is found in the β phase
The experimental results of the initial deposition experiments are shown in Table I. All five runs were performed at the same conditions, namely P = 500 W, m• MTS = 0.11 g/min, and QAr = 100 sccm. The run times varied randomly, as the plasma extinguished at unpredictable times. The efficiency is an indication of the percentage of MTS mass converted and deposited onto the quartz tubes. The layer thickness was calculated using deposition rate, run time, and assumed SiC density of 3.21 g/cm3 (Harris, 1995). SEM images of deposited SiC layers are shown in Figure 11 and Figure 12. These images were taken from experiment number B, and show what appear to be layered structures. Figure 12 shows agglomerates and smaller structures down to approximately 50 nm, supporting the presence of nanoparticles. The Journal of The Southern African Institute of Mining and Metallurgy
Figure 13 shows the XRD analysis of the SiC layers. No SiC peaks could be detected, indicating the presence of amorphous SiC. The absence of peaks could also be attributed to the small layer thickness and curvature of the quartz substrates. The small peak broadening at 25° is believed to be that of the amorphous quartz substrate. Typical EDAX results are shown in Table II, indicating the relative amounts of carbon, oxygen, silicon, and chlorine found in the SiC layers. The presence of oxygen and chlorine is indicative of SiO2 and HCl.
Conclusions and recommendations A microwave-induced plasma operating at atmospheric pressure was used to decompose methyltrichlorosilane to form SiC nanoparticle agglomerates with sizes down to 20 nm as determined from SEM and TEM micrographs. VOLUME 115
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Deposition of SiC at 15 kPa
Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma Table I
Experimental results of SiC deposition Exp no.
Run time (s)
Deposition rate (g/min)
Efficiency (%)
Layer thickness (μm)
A B C D E
320 900 360 160 180
0.013 0.009 0.008 0.018 0.016
12.16 18.48 9.692 9.898 12.08
8.352 15.12 5.927 5.854 5.776
Figure 11 – SEM image of SiC layers
Figure 12 – SEM image of SiC layer showing smaller nanostructures
Figure 13 – XRD analysis of SiC layers, indicating an amorphous structure
Table II
EDAX results indicating the relative elemental abundances Element
Weight %
Atomic %
Net int.
Net int. error
C O Si Cl
12.16 25.22 51.65 10.97
21.37 33.27 38.82 6.53
6.55 74.05 638.18 77.49
0.1 0.02 0.01 0.03
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Agglomerates were a common occurrence, resulting in larger particle size distributions when measured by the Zetasizer. The presence of β-SiC and silicon was confirmed using X-ray diffraction studies as well as optical spectroscopy. The first part of the investigation focused on the effect on particle size by varying the H2:MTS molar ratio and the total enthalpy. SEM, TEM, and Zetasizer results showed that higher enthalpy values and higher H2:MTS ratios produced smaller particle sizes. Furthermore, RSA results indicated that at high The Journal of The Southern African Institute of Mining and Metallurgy
Synthesis and deposition of silicon carbide nanopowders in a microwave-induced plasma
Acknowledgements The authors acknowledge the South African National Research Foundation for financial support, and the South African Nuclear Energy Corporation for use of their equipment.
References AHMED, Y.M.Z. and EL-SHEIKH, S.M. 2009. Influence of the pH on the morphology of sol–gel-derived nanostructured SiC. Journal of the American Ceramic Society, vol. 92. pp. 2724–2730. CARDOSO, R.P., BELMONTE, T., NOËL, C., KOSIOR, F., and HENRION, G. 2009. Filamentation in argon microwave plasma at atmospheric pressure. Journal of Applied Physics, vol. 105. pp. 093306-093306-8. DENG, J., SU, K., WANG, X., ZENG, Q., CHENG, L., XU, Y., and ZHANG, L. 2009. Thermodynamics of the gas-phase reactions in chemical vapor deposition of silicon carbide with methyltrichlorosilane precursor. Theoretical Chemistry Accounts, vol. 122. pp. 1–22. DHAGE, S., LEE, H.-C., HASSAN, M.S., AKHTAR, M.S., KIM, C.-Y., SOHN, J.M., KIM, K.-J., SHIN, H.-S., and YANG, O.B. 2009. Formation of SiC nanowhiskers by carbothermic reduction of silica with activated carbon. Materials Letters, vol. 63. pp. 174–176. HARRIS, G. 1995. Properties of Silicon Carbide, INSPEC, London. HONDA, S.-I., BAEK, Y.-G., IKUNO, T., KOHARA, H., KATAYAMA, M., OURA, K., and HIRAO, T. 2003. SiC nanofibers grown by high power microwave plasma chemical vapor deposition. Applied Surface Science, vol. 212–213. pp. 378–382. ISFORT, P., PENZKOFER, T., PFAFF, E., BRUNERS, P., GUNTHER, R.W., SCHMITZ-RODE, T., and MAHNKEN, A.H. 2011. Silicon carbide as a heat-enhancing agent in microwave ablation: in vitro experiments. Cardiovasculat Intervent Radiology, vol. 34. pp. 833–8. KAMLAG, Y., GOOSSENS, A., COLBECK, I., and SCHOONMAN, J. 2001. Laser CVD of cubic SiC nanocrystals. Applied Surface Science, vol. 184. pp. 118–122. KANEKO, T., MIYAKAWA, N., SONE, H., and YAMAZAKI, H. 2002. Growth kinetics of hydrogenated amorphous silicon carbide films by RF plasma-enhanced CVD using two kinds of source materials. Thin Solid Films, vol. 409. pp. 74–77. KÁROLY, Z., MOHAI, I., KLÉBERT, S., KESZLER, A., SAJÓ, I.E., and SZÉPVÖLGYI, J. 2011. Synthesis of SiC powder by RF plasma technique. Powder Technology, vol. 214. pp. 300-305. KATOH, Y., SNEAD, L., SZLUFARSKA, I., and WEBER, W. 2012. Radiation effects in SiC for nuclear structural applications. Current Opinion in Solid State and Materials Science, vol. 16. pp. 143–152. The Journal of The Southern African Institute of Mining and Metallurgy
KRAMIDA, A., RALCHENKO, Y., READER, J., and NIST-ASD-TEAM. 2014. NIST Atomic Spectra Database (version 5.2). National Institute of Standards and Technology, Gaithersburg, MD. http://physics.nist.gov/asd [Accessed 23 October 2014]. MOSHTAGHIOUN, B.M., POYATO, R., CUMBRERA, F.L., DE BERNARDI-MARTIN, S., MONSHI, A., ABBASI, M.H., KARIMZADEH, F., and DOMINGUEZ-RODRIGUEZ, A. 2012. Rapid carbothermic synthesis of silicon carbide nano powders by using microwave heating. Journal of the European Ceramic Society, vol. 32. pp. 1787–1794. PAPASOULIOTIS, G.D. and SOTIRCHOS, S.V. 1999. Experimental study of atmospheric pressure chemical vapor deposition of silicon carbide from methyltrichlorosilane. Journal of Materials Research, vol. 14. pp. 3397–3409. SACHDEV, H. and SCHEID, P. 2001. Formation of silicon carbide and silicon carbonitride by RF-plasma CVD. Diamond and Related Materials, vol. 10. pp. 1160–1164. SADDOW, S. and AGARWAL, A. 2004. Advances in Silicon Carbide Processing and Applications. Artech House, London. SATAPATHY, L.N., RAMESH, P.D., AGRAWAL, D., and ROY, R. 2005. Microwave synthesis of phase-pure, fine silicon carbide powder. Materials Research Bulletin, vol. 40. pp. 1871–1882. SAUTER, C., EMIN, M.A., SCHUCHMANN, H.P., and TAVMAN, S. 2008. Influence of hydrostatic pressure and sound amplitude on the ultrasound induced dispersion and de-agglomeration of nanoparticles. Ultrasonics Sonochemistry, vol. 15. pp. 517–523. SONE, H., KANEKO, T., and MIYAKAWA, N. 2000. In situ measurements and growth kinetics of silicon carbide chemical vapor deposition from methyltrichlorosilane. Journal of Crystal Growth, vol. 219. pp. 245–252. TANG, C.-J., FU, L.-S., FERNANDES, A.J.S., SOARES, M.J., CABRAL, G., NEVES, A.J., and GRÁCIO, J. 2008. Simultaneous formation of silicon carbide and diamond on Si substrates by microwave plasma assisted chemical vapor deposition. New Carbon Materials, vol. 23. pp. 250–258. TENDERO, C., TIXIER, C., TRISTANT, P., DESMAISON, J., and LEPRINCE, P. 2006. Atmospheric pressure plasmas: A review. Spectrochimica Acta Part B: Atomic Spectroscopy, vol. 61. pp. 2–30. TRUSOV, B. 2006. Terra - Phase and Chemical Equilibrium of Multicomponent Systems. Bauman Moscow State Technical University, Moscow. VAN LAAR, J.H., SLABBER, J.F.M., MEYER, J.P., VAN DER WALT, I.J., PUTS, G.J., and CROUSE, P.L. 2015. Microwave-plasma synthesis of nano-sized silicon carbide at atmospheric pressure. Ceramics International, vol. 41. pp. 4326–4333. VAßEN, R. and STÖVER, D. 2001. Processing and properties of nanophase nonoxide ceramics. Materials Science and Engineering: A, vol. 301. pp. 59–68. VENNEKAMP, M., BAUER, I., GROH, M., SPERLING, E., UEBERLEIN, S., MYNDYK, M., MÄDER, G., and KASKEL, S. 2011. Formation of SiC nanoparticles in an atmospheric microwave plasma. Beilstein Journal of Nanotechnology, vol. 2. pp. 665–673. WANG, X., SU, K., DENG, J., LIU, Y., WANG, Y., ZENG, Q., CHENG, L., and ZHANG, L. 2011. Initial decomposition of methyltrichlorosilane in the chemical vapor deposition of silicon-carbide. Computational and Theoretical Chemistry, 9 vol. 67. pp. 265–272. ◆ VOLUME 115
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enthalpy values (195–220 MJ/kg), particle size increases with increasing H2:MTS ratio. This trend could possibly be attributed to the increasing energy demand for hydrogen dissociation with increasing hydrogen amounts. The second part of the investigation included initial experiments on SiC layer deposition. Operating and synthesis conditions remained similar to those in the first part of the study, except the process was run at low pressure (15 kPa). Initial results indicate the successful deposition of SiC layers with thicknesses ranging between 5.8 and 15 µm. XRD results indicate amorphous crystalline structures, although further analysis is needed. The use of microwave-induced plasma shows promise for the deposition of SiC layers. Further experimental work is needed, however, in order to determine the quality and structure of these deposited layers. Pending these results, it is recommended that the use of microwave plasma systems be considered for the encapsulation of nuclear waste.
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a9
A redetermination of the structure of tetraethylammonium meroxidotrichlorido(thenoyltrifluoroacetyl acetonato-κ2-O,O')niobate(V) by R. Koen, A. Roodt and H. Visser
The tetraethylammonium salt of the mono-anionic coordination compound mer-oxidotrichlorido(thenoyltrifluoroacetylacetonato-κ2O,O’)niobate(V) (NEt4)[NbOCl3](ttfa)], has been prepared under aerobic conditions and characterized by single-crystal X-ray diffraction. (NEt4)[NbOCl3](ttfa)] crystallized in the monoclinic P21/c space group, with a = 11.483 (5), b = 12.563 (5), c = 17.110(5) Å, and β = 100.838 (5)º. The complex structure exists in a 50.0% (NbA) : 50.0% (NbB) positional disorder ratio. Keywords Bidentate, niobium(V), disorder.
Introduction Complexes containing organometallic type βdiketone ligands with O,O and O,N donor atoms are used widely in coordination chemistry and have applications in catalysis, radiopharmaceuticals, etc. (Schutte et al., 2011; Roodt, Visser and Brink, 2011; Brink et al., 2010; Otto et al., 1998). These ligand systems are very useful because of their highly coordinative nature, high solubility, and also due to their ability to be functionalized with various substituents on the carbonyl carbon atoms (Schutte et al., 2011). Only a small number of β-diketonate ligands have successfully been coordinated to a Nb(V) metal centre, with only a select few being characterized by X-ray crystallography (Viljoen, 2009; Bullen, Mason and Pauling, 1965; Preuss, Lamding and Mueller-Becker, 1994; Funk, 1934; Davies, Leedlam and Jones, 1999; Allen, 2002. The synthesis and crystal structure determination at room temperature of mer-oxidotrichlorido (thenoyltrifluoroacetylacetonato-κ2O,O’)niobate(V) {(NEt4)[NbOCl3(ttfa)]} was first reported by Daran et al. in 1979. Accordingly, for this current investigation, (NEt4)[NbOCl3(ttfa)] was re-evaluated at 100 K to determine if the structural features might change with temperature. This study of this structure forms part of an AMI-funded programme to better understand the solid-state characteristics of Ta(V) and Nb(V) complexes and the influences The Journal of The Southern African Institute of Mining and Metallurgy
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Experimental Materials and instruments All chemicals used for the synthesis and preparation of the complexes were of analytical grade and were purchased from SigmaAldrich, South Africa. The 1H-, 13C-, and 19F FT-NMR solutionstate spectra were recorded on a Bruker AVANCE II 600 MHz (1H: 600.28 MHz; 13C: 150.96 MHz; 19F: 564.83 MHz) or Bruker DPX 300 MHz (1H: 300.13 MHz; 13C: 75.47 MHz; 19F: 282.40 MHz) nuclear magnetic resonance spectrometer using the appropriate deuterated solvent. Chemical shifts, δ, are reported in ppm. 1H NMR spectra were referenced internally using residual protons in the deuterated solvents, Acetonitrile-d3 [CD3CN = 1.94(5) ppm]. 13C NMR spectra were similarly referenced internally to the solvent resonance [CD3CN = 1.39(4) ppm and 118.69(8) ppm] with values reported relative to tetramethylsilane (δ 0.0 ppm). The X-ray intensity data was collected on a Bruker X8 ApexII 4K Kappa CCD area detector diffractometer, equipped with a graphite monochromator and MoKα fine-focus sealed tube (λ = 0.71069 Å, T = 100(2) K) operated at 2.0 kW (50 kV, 40 mA). The initial unit cell determinations and data collection were done by the SMART (Bruker, 1998a) software package. The collected frames were integrated using a narrow-frame integration algorithm and reduced with the Bruker SAINT-Plus and XPREP software package (Bruker, 1999)
* Department of Chemistry, University of the Free State, South Africa. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. OCTOBER 2015
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Synopsis
of the electron-donating and -withdrawing groups of the β-diketone on the activity induced and reaction mechanisms at these metal centres.
Tetraethylammonium mer-oxidotrichlorido(thenoyltrifluoroacetylacetonato-κ2-O,O')niobate(V) respectively. Analysis of the data showed no significant decay during the data collection. The data was corrected for absorption effects using the multi-scan technique SADABS Bruker, 1998b) and the structure was solved by the direct methods package SIR97 (Altomare et al., 1999) and refined using the WinGX (Farrugia, 1999) software incorporating SHELXL (Sheldrick, 1997). The final anisotropic full-matrix least-squares refinement was done on F2. The aromatic protons were placed in geometrically idealized positions (C–H = 0.93 – 0.98 Å) and constrained to ride on their parent atoms with Uiso(H) = 1.2Ueq(C). Non-hydrogen atoms were refined with anisotropic displacement parameters. The graphics were obtained with the DIAMOND program (Brandenburg, 2006) with 50% probability ellipsoids for all non-hydrogen atoms.
Synthesis of (NEt4)[NbOCl3(ttfa)] (1) (Et4N)[NbCl6] (0.500 g, 1.147 mmol) was added to 4,4,4trifluoro-1(2-thienyl)-1,3-butanedione (ttfaH) (0.327 g, 1.147 mmol) in acetonitrile (20 cm3). The resulting solution was heated to 50ºC and stirred for 30 minutes. The excess solvent was evaporated and dark yellow plate-like crystals of the title compound (1), suitable for X-ray diffraction, were obtained (0.565 g, yield 89 %). IR (ATR, cm-1): ν(Nb=O) = 952. 1H NMR (300.13 MHz, Acetonitrile-d3, ppm): δ = 5.88 (s, 1H), 6.83 (m, 1H), 6.93 (dd, 1H), 7.40 (dd, 1H). 13C NMR (75.47 MHz, Acetonitrile-d3, ppm): δ = 30.1, 118.8, 123.9, 130.0, 137.4, 142.0, 182.3. 19F NMR (564.83 MHz, Acetonitrile-d3, ppm): -73.37.
Results and discussion The title compound was previously prepared by Daran et al. (1979), with X-ray diffraction data collected at room temperature. For this study, the reaction was modified as described above and the data collected at 100 K. The compound crystallizes in the monoclinic space group, P21/c, with four molecules in the unit cell (Z = 4). The asymmetric unit consists of a Nb(V) metal centre surrounded by three crystallographically independent chlorido groups (Cl1A – Cl3A), an oxido (O3A) and one O,O’-bonded thenoyltrifluoroacetonato ligand and a tetraethylammonium cation. A graphic illustration is shown in Figure 1. The complex molecule and the counter-ion are disordered over two positions in a 50 NbA: 50 NbB ratio as shown in Figure 2. General crystallographic details are presented in Table I, while selected bond lengths and bond angles are listed in Tables II and III respectively. A distorted octahedral geometry is displayed for NbA and NbB. The Nb-Claxial distances for NbA vary between 2.428(1) and 2.507(1) Å and the Nb1A-Cl1A and Nb1A-O3A have distances of 2.390(1) and 1.733(1) Å, respectively. When comparing Nb1A-O1A and Nb1A-O2A bond lengths, distances of 2.357(1) vs. 2.037(1) Å are observed. This difference is due to the effects of the electron withdrawing, CF3 substituent on the bidentate ligand backbone causing a longer NbA-O1A bond length. The trans Cl2-Nb1-Cl3 angle is 168.11(1)º, while the O1-Nb1-O2 bite angle is 79.65(1)º. A similar distortion is observed for NbB, with bond- lengths and angles in accordance with NbA. The coordination plane constructed through Cl1A, Cl2A, Cl3A, and O2A, as indicated in Figure 3, shows the niobium metal centre is slightly shifted out of this plane by 0.2751(3) Å. The molecular packing within the unit cell is illustrated in Figure 4. The packing illustrates a ‘head-to-tail’ arrangement
Table I
Crystallographic and refinement details of the title compound
Figure 1 – Graphic illustration of the mer-[NbOCl3(ttfa)] anion showing general numbering of atoms. Numbering of the disordered complex denoted by A = 50.0%. The displacement ellipsoids are drawn at 50% probability displacement level. Hydrogen atoms and counter-ion omitted for clarity
Crystallographic data
(NEt4)[NbOCl3(ttfa)]
Empirical formula Formula weight Crystal system, space group a, b, c (Å) α, β, γ (°) Volume (Å3), Z Density (calculated) (mg/m3) Crystal colour, crystal size (mm3) Absorption coefficient μ (mm-1) Theta range, F(000) Index ranges
C16H24C13F3N1Nb1O3S 566.68 Monoclinic, P21/c 11.483(5), 12.563(5), 17.110(5) 90.000, 100.838(5), 90.000 2424.3(16), 4 1.553 Yellow, 0.99 × 0.79 × 0.31 0.951 2.664 – 27.99°, 1144 -16<=h<=15, -15<=k<=15, -22<=l<=22 5834, 5209, 0.0574
Reflections collected, independent reflections, Rint Completeness (%) Max. and min. transmission Data, restraints, parameters Goodness-of-fit on F2 Final R indices [I>2sigma(I)]
Figure 2 – Graphic illustration of the mer-[NbOCl3(ttfa)] anion illustrating the disorder in an overlay. (red) NbA = 50.0%; (blue) NbB = 50.0%. Hydrogen atoms and counter-ion omitted for clarity
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R indices (all data) Largest diff. peak and hole (e.Å-3)
99.6 0.750 and 0.412 5834, 894, 501 1.0980 R1 = 0.0260 wR2 = 0.0740 R1 = 0.0303 wR2 = 0.0796 0.58, -0.58
The Journal of The Southern African Institute of Mining and Metallurgy
Tetraethylammonium mer-oxidotrichlorido(thenoyltrifluoroacetylacetonato-κ2-O,O')niobate(V) Table II
Selected bond lengths of the two disordered parts in the title compound mer-[NbOCl3(ttfa)] anion, denoted by NbA and NbB NbA
NbB
Atoms
Bond length (Å)
Atoms
Bond length (Å)
Nb1A-Cl1A Nb1A-Cl2A Nb1A-Cl3A Nb1A-O1A Nb1A-O2A Nb1A-O3A
2.390(1) 2.507(1) 2.428(1) 2.357(1) 2.037(1) 1.733(1)
Nb1B-Cl1B Nb1B-Cl2B Nb1B-Cl3B Nb1B-O1B Nb1B-O2B Nb1B-O3B
2.389(1) 2.373(1) 2.339(1) 2.254(1) 2.095(1) 1.745(1)
Table III
Selected bond angles of the two disordered parts in the title compound mer-[NbOCl3(ttfa)] anion, denoted by NbA and NbB NbA
NbB
Atoms
Bond angle (º)
Atoms
Bond angle (º)
O1A-Nb1A-O2A Cl1A-Nb1A-O3A O1A-Nb1A-O3A Cl1A-Nb1A-O2A Cl2A-Nb1A-Cl3A C2A-C3A-C4A
79.65(1) 96.28(1) 171.69(2) 163.58(1) 168.11(1) 120.60(2)
O1B-Nb1B-O2B Cl1B-Nb1B-O3B O1B-Nb1B-O3B Cl1B-Nb1B-O2B Cl2B-Nb1B-Cl3B C2B-C3B-C4B
76.82(1) 103.50(2) 166.36(2) 166.38(1) 160.12(1) 120.99(2)
Figure 4 – Packing of (NEt4)[NbOCl3(ttfa)] (NbA) along the c-axis. Displacement ellipsoids are drawn at the 50% probability level The Journal of The Southern African Institute of Mining and Metallurgy
Figure 5 – Illustration of the out-of-plane bend of the coordinated O,O'bonded thenoyltrifluoroacetonato ligand. Displacement ellipsoids are drawn at the 50% probability level VOLUME 115
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Figure 3 – Side view of the axial plane illustrating the out-of-plane distortion. Displacement ellipsoids are drawn at the 50% probability displacement level
when viewed along the c-axis. There are no classical hydrogen bonds or interactions observed in this structure. In Figure 5, two coordination planes are illustrated; the first one constructed through O1A, O2A, and Nb1A, and the second through O1A, C2A, C3A, C4A, and O2. The angle between planes revealed the slight, 1.677º out-of-plane bend of the coordinated O,O’-bonded thenoyltrifluoroacetonato ligand. This also contributes to the distorted octahedral geometry. The crystal structure determination of the published complex was performed at room temperature (298 K), while the synthesized analogue (1) was determined at 100(2) K. Data obtained for the title compound (1) correlates well with the previously published structure (Daran et al. 1979). The disordered part denoted by NbB differs less from the published structure and is probably a better representation of the anionic complex. Table IV illustrates a comparison between bond angles and distances of the published structure vs. the structure collected at 100K. The greatest difference
Tetraethylammonium mer-oxidotrichlorido(thenoyltrifluoroacetylacetonato-κ2-O,O')niobate(V) Table IV
Comparison of bond lengths and bond angles of (NEt4)[NbOCl3(ttfa)] (Bullen, Mason, and Pauling, 1965) collected at room temperature vs. (NEt4)[NbOCl3(ttfa)] at 100 K (NEt4)[NbOCl3(ttfa)]-NbA (100 K) Atoms
(NEt4)[NbOCl3(ttfa)] (298 K)
Bond length (Å)
Nb1-Cl1 Nb1-Cl2 Nb1-Cl3 Nb1-O1 Nb1-O2 Nb1-O3 Atoms O1-Nb1-O2 Cl2-Nb1-Cl3 C2-C3-C4
NbA
NbB
2.390(1) 2.507(1) 2.428(1) 2.357(1) 2.037(1) 1.733(1) Bond angle (º) NbA 79.65(1) 168.11(1) 120.60(2)
2.389(1) 2.373(1) 2.339(1) 2.254(1) 2.095(1) 1.745(1) NbB 76.82(1) 160.12(1) 120.99(2)
between the two complexes is the positional disorder observed in the newly synthesized product.
Bond length (Å)
Nb1-Cl1 Nb1-Cl2 Nb1-Cl3 Nb1-O1 Nb1-O2 Nb1-O3 Atoms
2.367(1) 2.365(2) 2.422(2) 2.285(3) 2.044(3) 1.704(3) Bond angle (º)
O1-Nb1-O2 Cl2-Nb1-Cl3 C2-C3-C4
78.7(1) 165.0(1) 122.9(1)
BRUKER AXS Inc. 1998a. Bruker SMART-NT Version 5.050, Area-Detector Software Package. Madison, WI, USA. BRUKER AXS Inc. 1998b. Bruker SADABS Version 2004/1, Area Detector Absorption Correction Software. Madison, WI, USA.
Conclusions A simplified method to obtain (NEt4)[NbOCl3(ttfa)] in aerobic conditions is reported. This highlights the misrepresentation of the ‘extreme sensitivity’ of niobium(V) complexes to air and water. Clearly, the exclusion of oxygen is not that important, while the exclusion of water is, since it will increase hydrolysis and thus the loss of chloride in favour of aqua, hydroxide, or oxo coordination. The crystallographic investigation revealed that this structure exhibited a 50:50 positional disorder. The electron withdrawing effects of the CF3 substituent on the bidentate ligand backbone is illustrated by the longer Nb-O bonds of Nb1A-O1A and Nb1B-O1B vs. Nb1A-O2A and Nb1B-O2B. All bond lengths and angles of the complex were found to be in accordance with similar structures in the literature.
Acknowledgements Financial assistance from the Advanced Metals Initiative (AMI) of the Department of Science and Technology (DST) of South Africa, through the New Metals Development Network (NMDN) managed by the South African Nuclear Energy Corporation Limited (Necsa) is gratefully acknowledged. Gratitude is also expressed towards SASOL, PETLabs Pharmaceuticals, and the University of the Free State for financial support of this research initiative outputs. Furthermore, this work is based on research supported in part by the National Research Foundation of South Africa (UIDs 71836 and 84913).
References ALLEN, F.H. 2002. Cambridge Structural Database (CSD) Version 5.35, November 2013 update. Acta Crystallographica, vol. B58. pp. 380–388. ALTOMARE, A., BURLA, M.C., CAMALLI, M., CASCARANO, G.L., GIACOVAZZO, C., GUAGLIARDI, A., MOLITERNI, A.G.G., POLIDORI, G., and SPAGNA, R. 1999. SIR97: a new tool for crystal structure determination and refinement Journal of Applied Crystallography, vol. 32. pp. 115-119. BRANDENBURG, K. 2006. DIAMOND, Release 3.0e. Crystal Impact GbR, Germany. BRINK, A., ROODT, A., STEYL, G., and VISSER, H.G. 2010. Steric vs. electronic anomaly observed from iodomethane oxidative addition to tertiary phosphine modified rhodium(I) acetylacetonato complexes following progressive phenyl replacement by cyclohexyl [PR3 = PPh3, PPh2Cy, PPhCy2 and PCy3]. Dalton Transactions, vol. 39. pp. 5572–5578.
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BRUKER AXS Inc. 1999. Bruker SAINT-Plus Version 6.02 (including XPREP),. Area-Detector Integration Software. Madison, WI, USA. BULLEN, G.J., MASON, R., and PAULING, P. 1965. The crystal and molecular structure of bis(acetylacetonato)nickel (II). Inorganic Chemistry, vol. 4. pp. 456–462. DARAN, J., JEANIN, Y., GUERCHAIS, J.E., and KERGOAT, R. 1979. The crystal structure of tetraethylammonium trichlorooxo(1,1,1-trifluoro-4-thenoyl2,4-butanedionato)niobate(V). Inorganica Chimica Acta, vol. 33. pp. 81–86. DAVIES, H.O., LEEDLAM, T.J., and JONES, A.C. 1999. Some tantalum(V) βdiketonate and tantalum(V) aminoalcoholate derivatives potentially important in the deposition of tantalum-containing materials. Polyhedron, vol. 18. pp. 3165–3174. FARRUGIA, L.J. 1999. WinGX suite for small-molecule single-crystal crystallography. Journal of Applied Crystallography, vol. 32. pp. 837–838. FUNK, H. 1934. Über die einwirkung von niob- und tantalpentachlorid auf organische verbindungen (IV. Mitteil.). Berichte der Deutschen Chemischen Gesellschaft, vol. 62. pp. 1801–1808. OTTO, S., ROODT, A., SWARTS, J.C., and ERASMUS, J.C. 1998. Electron density manipulation in rhodium(I) phosphine complexes: structure of acetylacetonatocarbonylferrocenyl diphenylphosphinerhodium(I). Polyhedron, vol. 17. pp. 2447–2453. PREUSS, F., LAMDING, G., and MUELLER-BECKER, S. 1994. Oxo- und thiotantantal(V)-verbindungen: Syntese von TaOX, und TaSX (X = OR, SR), Z. Zeitschrift für Anorganische und allgemeine Chemie, vol. 620. pp. 1812–1820. ROODT, A., VISSER, H.G., and BRINK, A. 2011. Structure/reactivity relationships and mechanisms from X-ray data and spectroscopic kinetic analysis. Crystallography Reviews, vol. 17. pp. 241–280. SCHUTTE, M., KEMP, G., VISSER, H.G., and ROODT, A. 2011. Tuning the reactivity in classic low-spin d(6) rhenium(I) tricarbonyl radiopharmaceutical synthon by selective bidentate ligand variation (L,L'-Bid, L,L ' = N,N', N,O & O,O' donor atom sets) in fac-[Re(CO)3(L,L'-Bid)(MeOH)]n complexes. Inorganic Chemistry, vol. 50. pp. 12486–12498. SHELDRICK, G.M. 1997. SHELXL97. Program for crystal structure refinement. University of Göttingen, Germany. VILJOEN, J.A. 2009. Speciation and interconversion mechanism of mixed halo O,O’- and N,O-bidentate ligand complexes of hafnium. MSc dissertation, Universtry of the Free State. 132 pp. ◆ The Journal of The Southern African Institute of Mining and Metallurgy
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A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form by C.J. Postma*, H.F. Niemand* and P.L. Crouse†
The separation of zirconium and hafnium is essential in the nuclear industry, since zirconium alloys for this application require hafnium concentrations of less than 100 ppm. The separation is, however, very difficult due to the numerous similarities in the chemical and physical properties of these two elements. Traditional methods for separation of zirconium and hafnium rely predominantly on wet chemical techniques, e.g. solvent extraction. In contrast to the traditional aqueous chloride systems, the AMI zirconium metal process developed by Necsa focuses on dry fluoride-based processes. Dry processes have the advantage of producing much less hazardous chemical waste. In the proposed AMI process, separation is effected by selective sublimation of the two tetrafluorides in an inert atmosphere under controlled conditions, and subsequent selective desublimation. Estimates are made for the sublimation rates of the two tetrafluorides based on the equilibrium vapour pressures. A sublimation model, based on the sublimation rates, was developed to determine if the concept of separation by sublimation and subsequent desublimation is theoretically possible. Keywords sublimation separation, zirconium tetrafluoride, hafnium tetrafluoride.
Introduction Zirconium requires several purification steps to conform to nuclear-grade specifications. Little information is available on the sublimation separation of Zr and Hf compounds, especially in the fluoride form, the majority of which deals with sublimation under vacuum conditions. On the industrial scale, only vacuum sublimation of ZrF4 has been reported. No records were found for the sublimation of ZrF4 in an inert atmosphere. Information on the sublimation rate of ZrF4 or HfF4 in an inert atmosphere is also limited. The rate is assumed to be dependent on several factors, of which temperature, area, and composition are considered the most important. In the process currently under investigation, the separation technique envisaged is by sublimation/desublimation in the tetrafluoride form. The separation involves the sublimation of the metal tetrafluorides in an inert atmosphere under controlled conditions, followed by desublimation (i.e. condensation) of the one metal fluoride in preference to the other. The Journal of The Southern African Institute of Mining and Metallurgy
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In this paper, a sublimation model is developed to predict the sublimation rates of both ZrF4 and HfF4 in an inert gas. These rates are used to calculate the partial pressures of the two fluorides exiting a sublimer and entering a desublimer where the one tetrafluoride is selectively removed from the gas stream, thereby separating the two tetrafluorides.
Literature survey and theoretical discussion Sublimation methods used for the separation of Zr and Hf are reported in literature, but these methods are all under vacuum conditions (Monnahela et al., 2013; Solov’ev and Malyutina, 2002a). Sublimation is, however, 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 et al., 2002; Pastor and Robinson, 1986; Solov’ev and Malyutina, 2002b; Yeatts and Rainey, 1965). The addition of baffles (Abate and Wilhelm, 1951; Kotsar’ et al., 2001; Yeatts and Rainey, 1965) is used quite frequently to help reduce the mechanical carry-over of impurities. These baffles are merely plates positioned between the charge and the cold finger. These impurities impart a greyish colour to ZrF4, whereas pure ZrF4 is much whiter. The literature also describes the use of a gettering agent (Monnahela et al., 2013; Solov’ev and Malyutina, 2002a), which seems to reduce the number of steps required to produce nuclear-grade ZrF4. Getters used include NiF2, zirconium oxides, and/or zirconium oxyfluorides.
* The South African Nuclear Energy Corporation SOC Ltd. (Necsa). † Department of Chemical Engineering, University of Pretoria. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Mar. 2015 and revised paper received July 2015. OCTOBER 2015
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Synopsis
A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form Area-dependent sublimation rate
Table I
MacFarlane et al. (2002) calculated the area-dependent rate of sublimation of ZrF4 and obtained a value of approximately 1.87 g/m2/s at 850–875°C.
Combined vapour pressure correlations for ZrF4
Product composition
Component
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 might be due to the accumulation of low-volatile components in the near-surface layer of the sample, making diffusion and evaporation increasingly difficult and resulting in decreased sublimation flux.
Layer height In a study on the influence of layer height on the vacuum sublimation rate of ZrF4, Ti, Esyutin, and Scherbinin (1990c) concluded that the sublimation rate does not necessarily depend on the height of the sample in the sublimator.
Vapour pressure of ZrF4 Figure 1 gives a range of vapour pressures from the literature for both ZrF4 and HfF4 at temperatures above 600°C (Benedict et al., 1981; Cantor et al., 1958; Koreneo et al., 1972; Sense et al., 1954, 1953). The data in Figure 1 was combined and can be expressed as Antoine constants for both ZrF4 and HfF4. These two expressions are given in Table I.
Experimental concept to be modelled The flow diagram for the proposed process is presented in
Figure 1 – Literature vapour pressures for ZrF4 (Benedict et al., 1981; Cantor et al., 1958; Koreneo et al., 1972; Sense et al., 1954, 1953)
and HfF4
ZrF4 HfF4
Vapour pressure [kPa] log(P) = 12.096−11879/T log(P) = 12.391−12649/T
Temperature range [°C] 600 – 920 600 – 950
Figure 2. The concept under consideration is to pass a stream of pre-heated dry nitrogen as a carrier gas over a bed of subliming ZrF4 and HfF4 (Figure 3); the gas exiting the sublimer then enters a desublimer operating at a slightly lower temperature. The difference between the vapour pressures as function of temperature is used to determine an optimum temperature for the desublimer. In the desublimer, one of the tetrafluorides is desublimed in preference to the other, thus effecting separation. The remainder of the gas mixture exits the desublimer and enters a water-cooled cold finger for collection of the remaining ZrF4 and HfF4, which can be subjected to a further separation step. The sublimer (Figure 3) consists of two rectangular sections; a bottom section containing the bulk mass to be sublimed, and a top section that facilitates the movement of the nitrogen gas and carries the sublimed tetrafluorides in the gas phase to the desublimer. The desublimer is a long cylindrical pipe that is heated to a predetermined temperature, depending on the partial pressures of the ZrF4 and HfF4 entering the desublimer.
Figure 3 – Sketch of the sublimer
Figure 2 – Block flow diagram for the sublimation separation of ZrF4 and HfF4
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A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form Modelling [5]
Flux model 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. As sublimation progresses, the bed height decreases with time, which causes changes in the mass transfer coefficient resulting in a change in the sublimation rate of the two tetrafluorides. The rate model for both ZrF4 and HfF4 is given in Equation [1]: [1] where ri is the sublimation rate of ZrF4 (or HfF4) in mol per unit sublimation area per unit time, ki,t is the mass transfer coefficient in m/s at time t, Pi* is the vapour pressure in kPa, pi’ is the partial pressure in the bulk gas, xi,t is the mol fraction of the respective tetrafluoride in the unsublimed bulk mass, R is the ideal gas constant (8.314 kPa.m3/kmol/K), and T is the temperature in K. In order to calculate the total flux along the length of the sublimation pan (nj,t), the pan is divided into segments each of length ΔL. 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 (Equation [2]). [2]
where Zt is the height of the headspace above the solid bed at any given time t.
Model parameters Mass transfer coefficient The mass transfer coefficient (ki) is required for calculating the rate of sublimation and is a function of the Sherwood number (Shi), the diffusion coefficient (DAB), and the equivalent flow diameter (De) (Equation [3]): [3] Sherwood numbers differ for each experimental set-up. In the case of convective mass transfer for forced convection over a flat plate (in this case a sublimation pan), and for laminar flow conditions with Reynolds number < 5 × 105, Prandtl number > 0.6, and Shmidt number (Sci) > 0.5, the Sherwood number can be calculated using Equation [4] (Çengel, 2006).
σAB is the collision diameter, a Lennard-Jones parameter in Å, where A refers to nitrogen and B to either ZrF4 or HfF4. Since σ 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 a HfF4 molecule assuming sphericity. The sizes of the respective molecules were calculated at room temperature with the use of SpartanTM software. 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 4.146 and 4.127 Å respectively. The collision integral (ΩD) is a dimensionless parameter and a function of the Boltzmann constant Κ (1.38 × 10-16 ergs/K), the temperature, and the energy of molecular interaction ∈AB. The boiling points (Tb) for ZrF4 (912°C) and HfF4 (970°C) (Lide, 2007) were used to calculate the values for ∈i with the use of an empirical correlation, given by Equation [6]: [6] Estimated values for the collision integrals for ZrF4 and HfF4 in N2 were calculated as 0.980 and 0.987, respectively. The diffusion coefficients were calculated at several temperatures and are listed in Table II.
Model results Limitations One practical problem encountered in the design of the experimental set-up is the working temperature of the valves (650°C). At this temperature, the vapour pressures are relatively low, which results in a very long sublimation time. Since all of the components of the valves are metallic, and it is general knowledge that these valves have a built-in safety factor, the decision was made to operate slightly above the maximum specification temperature of the valves, i.e. at 700°C.
Sublimation rates The initial load to be sublimed was taken as 80 g, which includes the HfF4 and other impurities. The full sublimation area is 0.0075 m2. The average area-dependant rates of sublimation at several temperatures calculated from the model are given in
[4]
Diffusion coefficients for ZrF4 and HfF4 in nitrogen
The diffusion coefficient can be estimated using the LennardJones potential to evaluate the influence of the molecular forces between the molecules. This correlation (Equation [5]), also known as the Chapman-Enskog equation, holds for binary gas mixtures of non-polar, non-reacting species (Green, 2008; Welty, 2001), which is the case for ZrF4 and HfF4 in nitrogen.
at a pressure of 1 atmosphere
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Temperature (°C)
(cm2/s)
(cm2/s)
700 750 800 850
0.693 0.757 0.824 0.893
0.677 0.740 0.806 0.874
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Table II
Diffusion coefficients
A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form Figure 4, which is a sum of the average rates of ZrF4 and HfF4 at several time intervals. It can be seen that the rate is a power function of the temperature, illustrating the effect of a higher temperature on the rate. The total sublimation time is dependent on the area of sublimation and the temperature of sublimation, the latter controls the two vapour pressures. The dependence of the total sublimation time on temperature is shown in Figure 5. At 850°C, the total sublimation time equals approximately 33 minutes, whereas at 700°C the total sublimation time was calculated to be approximately 24.5 hours. The vapour pressure ratio between 850 and 700°C is 42.7, which indicates that the rate at 700°C should be lower by at least this factor, which amounts to a total sublimation time of 23.5 hours. The other factor influencing the rate is the partial pressure in the gas stream, which is also higher at the higher temperature, which may contribute to the difference in the total sublimation time at 700°C. The mass sublimed with respect to time at 700°C is presented in Figure 6. Here it is evident that, according to the model calculations, some HfF4 will remain in the sublimation pan once all the ZrF4 has sublimed. Theoretically this implies that the sublimation can be stopped after a certain time, thereby separating most of the HfF4 from the ZrF4 in the cold finger. The rates of both HfF4 and ZrF4 sublimation are given in Figure 7. From this figure it is evident that the sublimation rate of HfF4 becomes increasingly significant as the sublimation progresses, i.e. as the bed height lowers with time. This is probably due to the mass fraction HfF4
increasing, since the ZrF4 has a higher rate of sublimation. The sublimed ZrF4 and HfF4 diffuse into the nitrogen stream, exit the sublimer, and enter a desublimer operating at a slightly lower temperature than that of sublimation. The desublimation temperatures for the two tetrafluorides can be calculated from the vapour pressure correlations based on the partial pressure of the respective fluorides entering the desublimer. Figure 8 indicates the desublimation temperatures obtained from the model at a sublimation temperature of 700°C for ZrF4 and HfF4. It is evident that a desublimer operating temperature of between 540°C (maximum temperature for HfF4) and 610°C (minimum temperature for ZrF4 desublimation) is required to ensure that, according to the model calculations, no HfF4 will desublime in the desublimer.
Figure 6 – Mass of ZrF4 and HfF4 sublimed with time at 700°C
Figure 4 – Rate of sublimation of impure ZrF4 at different temperatures
Figure 7 – Rate of sublimation for ZrF4 and HfF4 at 700°C as a function of the sublimation time
Figure 5 – Total sublimation time calculated at different sublimation temperatures
Figure 8 – Desublimation temperatures of ZrF4 and HfF4 at a sublimation temperature of 700°C
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A theoretical approach to the sublimation separation of zirconium and hafnium in the tetrafluoride form In the case of ZrF4, the desublimer operating temperature is still relatively high for all the ZrF4 to desublime while passing through the desublimer. A desublimer operating temperature of 30°C higher than that of the maximum temperature for HfF4 (i.e. 570°C) results in a vapour pressure of ZrF4 which is still relatively large, and some of the ZrF4 will therefore not desublime and will be collected on the cold -finger.
Comparison to literature data The area-dependent rates of sublimation for both ZrF4 and HfF4 at 850°C were calculated. The average rate over the entire duration of sublimation of the total sublimated mass amounts to approximately 5.36 g/m2/s, which is 2.87 times higher than the value estimated from literature data (1.87 g/m2/s). The difference between the literature and model rates may be attributable to the impurities present in the sample used by Macfarlane et al., (2002), since the presence of impurities can have a direct influence on the rate of sublimation.
Conclusions A sublimation model has been developed to predict the sublimation rates and the partial pressures of ZrF4 and HfF4 in the tetrafluoride form and in an inert gas. The gas exits a sublimer and enters a desublimer where the one tetrafluoride is desublimed in preference to the other, separating the two tetrafluorides. The model revealed an area-dependent sublimation rate that is 2.87 times higher than the value estimated from literature data (1.87 g/m2/s) at 850°C. This indicates that the rates obtained from the model are within an acceptable range. The difference between the literature and model rates may be attributable to the impurities present in the sample used by Macfarlane et al., (2002), since impurities can have a direct influence on the rate of sublimation. Due to experimental/equipment constraints, the operating temperature of the sublimer should be in the range of 700°C. Optimal temperature selection is imperative, since low temperatures result in a low sublimation rate and high temperatures increase the level of impurities in the sublimed product. At a sublimer operating temperature of 700°C, the model indicates that an operating temperature for the desublimer of between 540°C (maximum temperature for HfF4) and 610°C (minimum temperature for ZrF4 desublimation) is required to ensure that, according to the model calculations, only ZrF4 and no HfF4 will desublime. Selection of an optimal temperature of the desublimer is also critical, since too high a temperature will result in more ZrF4 lost to the cold finger. It is recommended that the model results be compared with experimental values, including if necessary the sublimation kinetics to account for any possible time dependencies of the vapour pressures.
References ABATE, L.J. and WILHELM, H.A. 1951. Sublimation of zirconium tetrafluoride. US Atomic Energy Commission, Ames Laboratory. BENEDICT, M., PIGFORD, T.H., and LEVI, H.W. 1981. Nuclear Chemical Engineering, McGraw-Hill, New York. ÇENGEL, Y.A. 2006. Heat and Mass Transfer. McGraw-Hill, New York. The Journal of The Southern African Institute of Mining and Metallurgy
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, no. 1-3. pp. 229–232. 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. GREEN, J.D. 2000. II / distillation / sublimation. Encyclopedia of Separation Science. 2nd edn). Wilson, I.D. (ed.). Academic Press, Oxford. KORENEO, Y., SOROKIN, I., CHIRINA, N., and NOVOSELO, A.V. 1972. Vapor-pressure of hafnium tetrafluoride. Journal of Inorganic Chemistry, vol. 17, no. 5. pp. 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. 2008. CRC Handbook of Chemistry and Physics. CRC, 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. AMI Precious Metals 2013 Conference, Protea Hotel, Cape Town, 14-16 October 2013. NEL, J.T., DU PLESSIS, W., CROUS, P.L., and RETIEF, W.L. 2011. Treatment of zirconia-based material with ammonium bi-fluoride. Patent WO2011013085 A1. PASTOR, R.C. and ROBINSON, M. 1986. Method for preparing ultra-pure zirconium and hafnium tetrafluorides. US patent 4,578,252 A. 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, no. 6. pp. 437–445. SOLOV’EV, A.I. and MALYUTINA, V.M. 2002a. Production of metallic zirconium tetrafluoride purified from hafnium to reactor purity. Russian Journal of Non-Ferrous Metals, vol. 43, no. 9. pp. 14–18. SOLOV’EV, A.I. and MALYUTINA, V.M. 2002b. 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. TI, V.A., ESYUTIN, V.S., and SCHERBININ, V.P. 1990a. 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. 1990b. The influence of the sample height on the zirconium tetrafluoride sublimation process. Kompleksn. Ispol’z. Miner. Syr’ya, vol. pp. 60–61. TI, V.A., ESYUTIN, V.S., and SCHERBININ, V.P. 1990c. The dependence of the zirconium tetrafluoride sublimation rate upon the process pressure. Kompleksnoe Ispol'zovanie Mineral'nogo Syr'ya, vol. 10. pp. 61–63. WELTY, J.R., WICKS, C.E., WILSON, R.E. and RORRER, G. 2001. Fundamentals of Momentum, Heat and Mass Transfer. Wiley, New York. YEATTS, L.B. and RAINEY, W.T. 1965. Purification of zirconium tetrafluoride. Technical report ORNL-TM-1292, US Atomic Energy Commission. ◆ VOLUME 115
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advanced metals initiative
Ferrous 2016 FERROUS AND BASE METALS DEVELOPMENT NETWORK CONFERENCE 2016 C o n f e r e n c e A n n o u n c e m e n t
15–17 October 2016 BACKGROUND Through its Advanced Metals Initiative (AMI) the South African Department of Science and Technology (DST) promotes research, development and innovation across the entire value chain of the advanced metals field. The goal of this initiative is to achieve sustainable local mineral beneficiation and to increase the downstream value addition of advanced metals in a sustainable manner. The achievement of this is envisioned to be through human capital development on post-graduate and post-doctoral level, technology transfer, localization and ultimately, commercialisation. The AMI comprises four networks, each focussing on a different group of metals. These are Light Metals, Precious Metals, Nuclear Materials and Ferrous and Base Metals (i.e. iron, steel, stainless steels, superalloys, copper, brass, etc.). The AMI FMDN 2015 Conference aims to bring together stakeholders from the mineral sector, academia, steel industry, international research institutions and government in order to share and debate the latest trends, research and innovations, specifically in the areas of energy, petrochemical, corrosion, materials for extreme environments and transport, local mineral beneficiation and advanced manufacturing related to these materials. Keynote speakers to be invited include international specialists in the fields of ferrous metals, computational materials science, high temperature corrosion and mineral beneficiation. The Ferrous and Base Metals Development Network (FMDN) of the DST’s Advanced Metals Initiative (AMI) programme will host the AMI’s annual conference in 2016. The conference seeks to share insight into the state of R&D under the AMI-FMDN programmes and explore and debate the following broad themes:
Development of high performance ferrous and base metal alloys for application in the energy and petrochemical industries Development of corrosion resistant ferrous and base metal alloys Development of lightweight and/or durable steels for cost-effective transportation and infrastructure, and Panel discussions on possible future value-adding R&D programmes under FMDN within the South African National Imperatives.
OBJECTIVES
WHO SHOULD ATTEND
Insight into ferrous and base metal materials R&D for application
Stakeholders from the energy, petrochemical, corrosion and transportation industries where ferrous (i.e. iron, steel, stainless steels, superalloys, etc.) and base (i.e. copper, brass, etc.) metals are used in their infrastructure. Also included in this invitation are local and international Higher Education Institutions (HEIs), Government Departments and Science Councils that are involved and/or have interest in R&D in these areas.
in the areas of energy, petrochemical, corrosion, extreme environments, improved processing technologies and advanced alloys for the transport industry in South Africa and globally.
EXHIBITION/SPONSORSHIP Sponsorship opportunities are available. Companies wishing to sponsor or exhibit should contact the Conference Co-ordinator.
our future through science
For further information contact: Head of Conferencing, Raymond van der Berg SAIMM, P O Box 61127, Marshalltown 2107 Tel: +27 11 834-1273/7 ·Fax: +27 11 833-8156 or +27 11 838-5923 E-mail: raymond@saimm.co.za · Website: http://www.saimm.co.za
http://dx.doi.org/10.17159/2411-9717/2015/v115n10a11
Glow discharge optical emission spectroscopy: a general overview with regard to nuclear materials by S.J. Lötter*†, W. Purcell* and J.T. Nel†
Synopsis Glow discharge optical emission spectroscopy (GD-OES) is an analytical technique used in the analysis of solid, conducting materials. Though primarily of interest as a depth profiling technique on samples with varying layers of both conducting and non-conducting materials, it is also capable of rapid bulk analysis of homogenous, solid samples. GD-OES combines the advantages of ICP-OES (wide detection range, speed, and lack of interferences) with the solid sampling of XRF techniques. This allows the analyst to not only quantify the elemental composition of a sample, but to evaluate it in terms of homogeneity with depth, a field in which auger electron spectroscopy (AES) and secondary ion mass spectrometry (SIMS) have traditionally been the primary techniques. Although GD-OES does not replace these useful techniques, it does offer various advantages over them, making it an excellent complementary analytical tool. In GD-OES analysis, a low-pressure glow discharge plasma is generated with the sample material acting as a cathode, accelerating the cations in the plasma towards the sample surface. This bombardment causes the sample material to ‘sputter’, essentially knocking free atoms or molecules of analyte material, which are then drawn into the plasma where they are excited. The light emitted by this excitation is then diffracted to separate the wavelengths emitted by the specific elements and detected by a spectrophotometer. The intensity of the signal is directly proportional to the quantity of analyte element present in the sample, allowing simple calibration and quantitative determination of the elements. Technological improvements made in the past twenty years or so have significantly increased the usefulness of GD-OES for surface analysis. Faster plasma stabilization and start-up allow for the quantification of surface layers as thin as 1 nm. Sputter rates have been accurately measured for many common materials, allowing them to be built into software libraries in the instrument’s control software. This dramatically expands the usefulness of this software and the ease of performing analyses. GD-OES analysis of nuclear materials allows for the rapid determination of the elemental composition without the requirement of initial dissolution. The thickness of any corrosion layers on nuclear materials can also be determined. Keywords glow discharge optical emission spectrometry, GD-OES, nuclear materials.
profiles of materials that contain layers of varying composition and thickness allows not only for the quantification of a sample, but also the evaluation of its composition in terms of homogeneity. Unfortunately the method is not yet as well established as other techniques. In order to be at its most effective it requires conductive samples. GD-OES was developed by Grimm in 1967 (Payling, Jones, and Bengtson, 1997). It was initially used to analyse solid metallic samples (bulk analysis) where all surface effects were eliminated by a pre-burn. In the 1970s, Roger Berneron, among others, began investigating the phenomena that occurred during this pre-burn period and developed GD-OES as a surface analytical technique. In 1988 M. Chevrier and Robert Passetemps first successfully applied an RF potential to a conventional Grimm-type source (GD) to produce an RF device that was able to also analyse non-conductive samples. Currently, GD-OES is still used mainly in the analysis of solid metallic samples (Azom.com, 2004). Its most useful application is as a depth profiling technique on samples with varying layers of both conducting and non-conducting materials (Shimizu et al., 2003), but it is also capable of rapid bulk analysis of homogenous samples (both conducting and non-conducting). A typical analysis takes only a few minutes, which is one of the main advantages of GD-OES, especially where quick results are required, for example at a smelter. In GD-OES analysis, when using a Grimm-type direct current (DC) lamp source, a potential difference of the order of 1 kV is applied between two electrodes in a cell containing a gas, usually argon, at low pressure (around 1 Torr) (Bogaerts and
Introduction
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* Department of Chemistry, University of the Free State, South Africa. † The South African Nuclear Energy Corporation SOC Ltd. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received July 2015 and revised paper received Aug. 2015. OCTOBER 2015
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Glow discharge optical emission spectroscopy (GD-OES) combines the detection range, speed, and lack of interferences of an inductively coupled plasma optical emission spectroscopy (ICP-OES) instrument with solid sampling reminiscent of X-ray fluorescence (XRF) techniques. Its ability to perform depth
Glow discharge optical emission spectroscopy Gijbels, 1998). This causes the gas to ionize into positive and negative ions, forming the plasma. The electrons are accelerated towards the anode, sustaining the plasma and generating more positive ions. The positive ions are similarly accelerated towards the cathode, the sample material. The resulting bombardment of the sample material causes it to ’sputter’, essentially knocking free atoms or molecules of analyte material. These atoms are then drawn into the plasma where they too are excited. A diagram of this process can be seen in Figure 1. A radio frequency (RF) lamp may also be used for analysing non-conducting samples (Winchester and Payling, 2004), see Figure 2. The sputtered analyte is then drawn into the high-energy plasma where it is excited into a higher electronic energy state. In order to return to their ground state the analyte ions emit a photoelectron of a wavelength characteristic of the element emitting it. In GD-OES the characteristic light emitted by this excitation/de-excitation process is passed into a spectrophotometer through an entrance slit, and is then diffracted by a concave holographic mirror grating that contains 3600 lines per millimetre. This is done in order to separate the emitted light into wavelengths that are specific for each element. The
Figure 1 – Schematic of the various processes occurring during glow discharge (GD) (Betti and de las Heras, 2004)
Figure 2 – Diagram of a Rowland circle used in the LECO GDS850A (Leco Corporation, 2007)
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individual wavelength intensities can then be detected using either a charge-coupled device (CCD) or a photomultiplier tube (PMT). The intensity of the signal is directly proportional to the quantity of analyte element present in the plasma, allowing simple calibration and quantitative determination of elements. In practice this is often achieved using a Rowland circle (Figure 3) with PMTs positioned behind exit slits that are positioned at the wavelength corresponding to characteristic emission lines of various elements. This allows the instrument to read all of these lines simultaneously, greatly increasing sample throughput rate at the cost of the ability to generate a complete spectrum. As there is only a single PMT installed for most elements in this type of spectrometer, detection of interferences is more difficult than with scanning sequential instruments. A sequential instrument would, however, not be capable of the most useful aspect of the GD-OES, which is its ability to perform depth profiling. The signals that are captured by the PMTs are recorded by a computer at a typical rate of 10 000 to 30 000 measurements per second. Technological improvements made in the past twenty years or so have significantly increased the usefulness of GDOES for surface analysis (Azom.com, 2004). Faster plasma stabilization and start-up times allow for the quantification of surface layers as thin as 1 nm. Sputter rates have been accurately determined for many common materials, allowing them to be built into the software libraries of the instrument’s control software. This significantly expands the usefulness of the software. Improvements in calibration systems allow for the changing of instrumental conditions without the need to completely recalibrate, saving time and effort. In glow discharge mass spectrometry (GD-MS), a mass spectrometer is attached to the GD source rather than an optical detector. This allows for a much lower limit of detection at the cost of linear dynamic range. The instrument is otherwise essentially the same as a GD-OES instrument. A collision cell can be used at the interface between the glow discharge plasma and mass spectrometer in order to minimize interference from polyatomic species formed in the plasma.
Figure 3 – Conceptual diagram of a radio frequency glow discharge device (RF GD) (Winchester and Payling, 2004) The Journal of The Southern African Institute of Mining and Metallurgy
Glow discharge optical emission spectroscopy As GD-MS still makes use of a glow discharge plasma sampling system it shares the GD-specific advantages and limitations of the GD-OES. Due to the cost of a mass spectrometer a GD-MS instrument is generally significantly more expensive than an OES instrument. Modern instruments are equipped with sophisticated software packages that use libraries containing the sputter rates of known materials and advanced algorithms. These are used to calculate both the depth and mass composition of the layers, as seen in Figure 4 and Figure 5. Figure 4 shows the depth profile of a hard-disk platter with seven distinct layers clearly visible, while Figure 5 shows the composition of the surface of a photovoltaic cell.
Figure 4 – Three replicate quantitative depth profiles (QDPs) of the surface of a hard disk exhibiting seven resolved layers at a depth of 100 nm (Leco Corporation, 2007)
Figure 6 – Cross-section of a glow discharge lamp showing vacuum, cooling, and gas flow lines (Payling, Jones, and Bengtson, 1997)
Figure 5 – Depth profile of the absorber layer of a copper indium gallium selenide (CIGS) thin film photovoltaic (PV) cell by pulsed RF GD-OES (Horiba Scientific, 2014)
Figure 7 –The LECO GDS850A (Leco Corporation, 2007)
Instrumentation
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A GD-OES instrument is a fairly simple piece of equipment to operate, in comparison to other instruments capable of similar analyses. Once a sample is properly prepared, sample introduction consists of simply placing a sample over the Oring at the lamp opening and allowing the vacuum to hold it in place. A simple cross-section of a standard Grimm-type GD lamp can be seen in Figure 6.
When positioning a sample it is preferable to ensure that no previous GD craters overlap with the surface to be analysed. Care must be taken that samples are non-porous and smooth. If they are not, the vacuum will not be strong enough to form the GD plasma, the instrument will detect this, and the sample will be ejected before analysis can begin. If the seal is tight the sample will then be held in place by the reamer, a pneumatically controlled ‘air jack’ with an integrated drill bit. This reamer serves multiple functions, including holding the sample in place, cleaning the anode between analyses, and serving as a conductor to allow the sample to act as cathode (when using the DC lamp). When moving into position the reamer impacts the sample with some force, which can break brittle samples. The LECO GDS850A (Figure 7) supports both a DC and an RF lamp. The RF lamp is intended for use with nonconducting materials and has a lower effective applied power. It was found that using this lamp with the copper anode allowed only for the analysis of very thin, regular, nonconducting layers on a conducting substrate. The DC lamp is able to generate a greater applied power, which can increase sensitivity but is also more prone to short-circuits.
Glow discharge optical emission spectroscopy In order to ensure that the Rowland circle is in alignment, the instrument is equipped with two PMTs mounted at diffraction angles for detecting the characteristic wavelengths of iron. These two iron ‘lines’ are used when profiling the instrument. This is done by analysing a sample containing iron. Detecting these two lines allows the instrument to confirm the position of the spectrometer’s entry slits. As the instrument stores calibration data it is necessary to driftcorrect a method’s calibration before quantitative analysis. This is performed by analysing standard samples, preselected in the method, and adjusting the calibration according to the intensities detected. This drift must be done whenever a different method is used but the warm-up and profiling need to be performed only once per working period.
The use of GD-OES in the nuclear industry As previously mentioned, the most striking capability of GDOES is its ability to accurately and rapidly determine a depth profile of multiple layers of varying composition on a substrate. The raw data from a GD-OES analysis is given in time versus intensity of emission, as can be seen in Figure 8, where layers of UO2F2·1.5H2O, and UF4 are seen bonded to a nickel substrate. This data, along with known sputter rates for each material, allows for the characterization of the various uranium species present on the surface of the material. In the 1980s the erstwhile Atomic Energy Corporation of South Africa (AEC) used GD-OES intensively to study surface and interface phenomena like corrosion, surface cleanness, passivation of surfaces, decontamination efficiency, plating, etc. (Nel, 1991). Valuable information could be obtained by analysing surfaces of materials that were exposed to UF6, as indicated in Figure 8. The types of uranium compounds that were found on a surface were used by nuclear engineers and scientists to choose the best materials for construction or to diagnose chemical and material problems that occurred during plant operation. GD-OES was also routinely used for the bulk analysis of materials that are used in nuclear and uranium enrichment
plants. Examples of these are high nickel-containing alloys, normal and special stainless steels (e.g. 316, 304, and 17PH), Monel, Inconel, aluminium alloys, copper, and brass, etc. The impurity specifications for nuclear materials are usually extremely strict and the low limit of detection of GDOES makes this technique extremely powerful in the nuclear industry. In the current programme of the Advanced Metals Initiative (AMI), and especially the Nuclear Materials Development Network (NMDN), one of the main focuses is the development of nuclear-grade zirconium metal and the study of corrosion of zirconium alloys. GD-MS has been used in the analysis of spent uranium oxide fuel (Robinson and Hall, 1987) prior to recycling. It is necessary to confirm that the reprocessing plant is working efficiently and that fission products, activation products, and corrosion products are successfully removed. GD-MS offers a faster and more quantitative approach than traditional methods, such as spark source mass spectrometry (SSMS), without the need for time-consuming interpretation of results by highly skilled operators.
Experimental Samples of Zircaloy 2 and Zircaloy 4, obtained from ATI Wah Chang, were exposed to air at 600°C for periods of either 16 or 32 hours. The samples were allowed to cool rapidly in air and then analysed by GD-OES to determine the thickness of the oxide layer formed. A sample of zirconium metal, obtained from Alfa Aesar, was analysed as well in order to provide a baseline from which to compare the degree of the growth of the oxide layer. The elemental composition of the Zircaloy materials can be seen in Table I. The GD-OES instrument was calibrated using NIST standards 1212a, 1234, 1235, 1236, 1237, 1238, and 1239, obtained from Pelindaba Analytical Laboratories. The photomultiplier tube power settings were determined using the automated software tool built into the control software, with applied voltages ranging between 400 and 900 V. Instrument settings can be seen in Table II. As the RF lamp was used the Teflon isolation puck was utilized in all determinations and was kept at a constant temperature of 12°C by an external chiller. This puck is similar to the cooling puck used with the DC lamp but, in addition to cooling, also electrically isolates the sample from the reamer.
Results and discussion Oxygen readings in all cases were higher than expected, but this was likely a calibration issue. As the oxygen calibration values were for dissolved oxygen in the metal standards rather than the stoichiometric levels found in the oxide
Table I
Elemental composition of Zircaloy 2 and Zircaloy 4 (Roskill Information Services, 2011, p. 322)
Figure 8 – Depth profile of a thin layer containing UO2F2·1.5H2O and UF4 on a nickel substrate (Nel, 1991)
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Sample
Zr %
Sn %
Fe %
Ni %
Cr %
Zircaloy 2 Zircaloy 4
98.23 98.28
1.5 1.5
0.12 0.12
0.05 -
0.1 0.1
The Journal of The Southern African Institute of Mining and Metallurgy
Glow discharge optical emission spectroscopy Table II
GDS850A settings Setting
Value
Voltage (V) True plasma power (W) Vacuum (Torr) Analysis time (min)
1200 20 6 - 12 10
Figure 12 – Quantitative depth profiles of Zircaloy 2 after 16 and 32 hours’ exposure in air at 600°C
Figure 9 – Quantitative depth profile of a clean plate of zirconium metal before corrosion
Figure 13 – Quantitative depth profiles of Zircaloy 4 after 16 and 32 hours’ exposure in air at 600°C
Table III
Measured thicknesses of Zircaloy corrosion layers
Figure 11 – Quantitative depth profiles of Zircaloy 2 and 4 after 32 hours’ exposure in air at 600°C The Journal of The Southern African Institute of Mining and Metallurgy
Sample
Time at 600°C (h)
Layer thickness (μm)
Zircaloy 2 Zircaloy 2 Zircaloy 4 Zircaloy 4
16 32 16 32
35 100 31 60
layers, the intensities detected were several orders of magnitude higher than the calibration range. Oxygen values must thus be described as semi-quantitative. As the objective of the study was to determine the thickness of the layers rather than their elemental composition, this was not considered too great of a difficulty. The oxide layer formed naturally at room temperature by zirconium metal was found to be approximately 100 nm thick with approximately another 200 nm required for the transition to mostly pure zirconium, as seen in Figure 9. The thickness of the corrosion layers varied between 31 and 100 µm, depending on the duration of and temperature at which the corrosion test was performed (Table III). The least VOLUME 115
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Figure 10 – Quantitative depth profiles of Zircaloy 2 and 4 after 16 hours’ exposure in air at 600°C
Glow discharge optical emission spectroscopy corrosion was exhibited by Zircaloy 4, with layer thickness marginally thinner than those seen with Zircaloy 2 at 16 hours but significantly thinner after both had been corroded for 32 hours. These results are clearly visible in Figure 10 and Figure 11. The increase in oxide layer thickness of 285.7 % for Zircaloy 2 can clearly be seen in Figure 12, while the Zircaloy 4 showed an increase of only 193.5 % (Figure 13). As expected, Zircaloy 4 exhibited a significantly better resistance to corrosion than Zircaloy 2. These results clearly show the obvious utility of GD-OES for the rapid determination of oxide layers on nuclear fuel cladding. It can be highly useful as a method to quantify the corrosion of known and experimental materials for nuclear application. Its usefulness extends to the determination of layer thickness and composition for known and potential coatings on nuclear and non-nuclear materials.
References
BOGAERTS, A. and GIJBELS, R. 1998. Fundamental aspects and applications of glow discharge spectrometric techniques. Spectrochimica Acta Part B, vol. 53. pp. 1–42. HORIBA SCIENTIFIC. 2014. Ultra fast elemental depth profiling. http://www.horiba.com/fileadmin/uploads/Scientific/Documents/Emission /GDPROFILER_Series.pdf [Accessed 18 April 2015]. LECO CORPORATION. 2007. GDS850A Glow Discharge Spectrometer. http://www.leco.com/component/edocman/?task=document.viewdoc&id=4 9&Itemid=0 [Accessed 4 September 2013]. NEL, J.T. 1991. Oppervlakanalises van uranielfluoriedlagies met 'n gloeiontladingslamp. DPhil thesis, University of Pretoria. PAYLING, R., JONES, D. and BENGTSON, A. 1997. Glow Discharge Optical Emission Spectrometry. Wiley, Hoboken, NJ. ROBINSON, K. and HALL, E.F.H. 1987. Glow discharge mass spectrometry for nuclear materials. Journal of Metals, vol. 39. pp. 14–16. ROSKILL INFORMATION SERVICES. 2011. Zirconium: Global Industry Markets and Outlook. 13th edn. London. SHIMIZU, K., HABAZAKI, H., SKELDON, P. and THOMPSON, G.E. 2003. Impact of RFGD-OES in practical surface analysis. Spectrochimica Acta Part B, vol. 58. pp. 1573–1583.
AZOM.COM. The A to Z of Materials. 2004. http://www.azom.com/article.aspx?ArticleID=2449 [Accessed 9 July 2013]. BETTI, M. and DE LAS HERAS, L.A. 2004. Glow discharge spectrometry for the characterization of nuclear and radioactively contaminated environmental samples. Spectrochimica Acta Part B, vol. 59. pp. 1359–1376.
WINCHESTER, R.M. and PAYLING, R. 2004. Radio-frequency glow discharge spectrometry: a critical review. Spectrochimica Acta Part B, vol. 59. pp. 607–666. ◆
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a12
The influence of niobium content on austenite grain growth in microalloyed steels by K.A. Annan, C.W. Siyasiya and W.E. Stumpf
The relationship between niobium content and austenite grain growth has been investigated through hot rolling simulation on a Bähr dilatometer. The effect of delay time between passes during rough rolling in Nbmicroalloyed steels with nitrogen contents typical for electric arc furnace (EAF) melting was studied. The results indicate that the grain growth constants n, Q, and A increase with an increase in Nb content. The activation energy for austenite grain growth Q was found to be in the range of 239 to 572 kJ/mol, the exponential constant n ranged from 2.8 to 6.2, and the material and processing condition constant A from 4.24 × 1012 to 4.96 × 1028, for steels with niobium contents ranging from 0.002% Nb to 0.1% Nb. A general constitutive equation for the prediction of austenite grain growth in these Nb-microalloyed steels under rough rolling conditions has been developed. Good agreement between the experimental and the predicted values was achieved with this constitutive equation. Keywords constitutive equation, austenite grain growth, microalloying, deformation.
Introduction Grain refinement has been found to increase both the strength and toughness of steels (Gao and Baker, 1998; Sharma, Lakshmanan, and Kirkaldy, 1984; Seok et al., 2014; Maalekian et al., 2012). It is also known that the austenite grain size directly influences the microstructure, and thus the mechanical properties, of the steel (Maalekian et al., 2012; Yue et al., 2010). Effective grain growth control is reported to be achieved through addition of precipitate-forming elements, such as Nb that, slow down the grain boundary migration through pinning and solute drag mechanisms (Yu et al., 2010; Olasolo et al., 2011; Alogab et al., 2007). Much work has been carried out on austenite grain lsize control by the addition of precipitate-forming elements that have a strong affinity for interstitial elements, such as carbon and nitrogen, which form the dispersed pinning particles to inhibit the austenite grain growth (Alogab et al. 2007; Hodgson and Gibbs, 1992; Nanba et al., 2003; Flores and Martinez, 1997; Rollett, Srolovitz and Anderson, 1989). Austenite grain growth can be described using a conventional or a modelling approach. With conventional approaches (e.g. metallography), The Journal of The Southern African Institute of Mining and Metallurgy
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Materials and techniques Table I shows the chemical compositions of the five microalloyed steels used in the study, which were cast by vacuum induction melting into ingots of 16 kg. The compositions were chosen to test the effect of the Nb content on the austenite grain growth while the other elements were kept approximately constant. Note, however, that the melting and casting procedure used for these laboratory steels
* Department of Materials Science and Metallurgical Engineering, University of Pretoria, South Africa. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015 and revised paper received Aug. 2015. OCTOBER 2015
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Synopsis
in-situ monitoring of austenite grain growth at high temperatures is impossible. Modelling the grain growth behaviour based on available data is the ideal option. To quantitatively describe austenite grain growth therefore requires development of a sound mathematical austenite grain growth equation that accounts for the effects of the varying microalloying elements in inhibiting austenite grain growth. Numerous attempts have been made to develop an empirical model based on the general equation developed by Sellars and Whiteman (1979). Many of these models do not account for the direct effects of the microalloying elements such as Nb in austenite grain growth control [Fu et al., 2011; Wang and Wang, 2008; Wang et al., 2006; Shanmugama et al., 2005; Banerjee et al., 2010; Pous-Romeroa et al., 2013). The current work has considered this limitation, taking into account the direct effect of niobium in grain growth control during thermal processing. This is done by incorporating the initial austenite grain size Do and the microalloying element niobium in the development of a constitutive equation for grain growth prediction in Nb-containing microalloyed steels.
The influence of niobium content on austenite grain growth in microalloyed steels Table I
Chemical composition of the steels used in the study Steel
0.002% Nb 0.03% Nb 0.05% Nb 0.07% Nb 0.10% Nb
C
Mn
Si
0.17 0.16 0.17 0.14 0.18
1.13 1.15 1.15 1.10 1.27
0.11 0.15 0.15 0.15 0.15
Chemical composition (wt %) Nb Ti V 0.002 0.03 0.05 0.07 0.10
resulted in somewhat higher nitrogen contents, which are nearer to those expected from electric arc furnace (EAF) than from basic oxygen furnace (BOF) steelmaking.
Compression test on a Bähr dilatometer Cylindrical samples (5 mm diameter × 10 mm length) machined from the laboratory cast slabs were heated by induction in the Bähr dilatometer to 1150°C at a rate of 5°C s-1 and held at 1150°C for 5 minutes to achieve homogenization. The samples were then cooled at 5°C s-1 to the testing temperature (1000, 1050, 1100, and 1150°C) where a singlehit compression was applied after holding at the test temperature for 20 seconds. A strain of 0.4 was used at a strain rate of 0.1 s-1. After compression, the samples were then held at the deformation temperature in the Bähr dilatometer for times of 0, 10, 30, 60, 90, or 120 minutes to simulate the delay time after rough rolling, before rapid cooling at a rate of 600°C s-1 to room temperature. Oxidation of the specimens during compression was prevented by passing a continuous flow of high-purity helium through the system. The cooled samples were then tempered at 490°C for 72 hours to improve the response of prior austenite grain boundaries to etching. In this study, the grain size of the samples held for zero (0) minutes after deformation was used as the initial grain size D0 for the steel at that temperature. The scheduled profile followed in the single hit compression tests is shown in Figure 1. The samples were mounted, ground, polished, and then etched with picric acid solution containing 100 ml saturated aqueous picric acid, 100 ml distilled water, 4 g sodium
0.001 0.02 0.02 0.02 0.02
0.007 0.008 0.009 0.006 0.009
Al
N
Ni
0.05 0.02 0.03 0.04 0.04
0.018 0.018 0.019 0.020 0.020
0.15 0.15 0.15 0.17 0.17
tridecylbenze sulphonate, and 2-3 drops of triton. Etching was carried out at a temperature of 60–70°C for times ranging from 3 to 14 minutes. The samples were observed under an Olympus BX51MTM optical microscope to reveal the prior austenite grain boundaries. The austenite grain size was measured using the average linear intercept method according to ASTM E112 (Van der Voort, 1984). To obtain a statistically acceptable grain size distribution, more than 300 intercepts were measured on each sample.
Results The solubility behavior of precipitates in microalloyed steels predicted by Thermo-CalcTM The volume fractions of precipitates as a function of the temperature predicted by Thermo-CalcTM for the studied steels are shown in Figure 2. These show that the reheating temperature range of 800–1250°C will lead to dissolution of a substantial amount of Ti,Nb(C,N) carbonitrides but not to a complete dissolution of all of these precipitates. NbN precipitates within the temperature region of 1000–1200°C, while TiN does not go into solution at temperatures considered in this study. The higher dissolution temperatures shown for NbN in these microalloyed steels are due to their higher N content of about 0.019 wt%, which is more typical for steel produced in an EAF (N in solute-rich regions preferentially combines with Nb). The presence of precipitates means that a unimodal grain structure can be predicted if these precipitates are effective in pinning the grain boundaries (Fernández, Illescas and Guilemany, 2007; Van der Voort, 1984; Zener, 1948 (as cited by Ringer, Li and Easterling, 1989). The volume fraction of Nb precipitates with respect to temperature as predicted by Thermo-CalcTM is shown in Figure 3. This shows that the volume fraction at a given temperature is highly dependent on the Nb content in the steel, thus the higher the Nb content, the higher the volume fraction of the precipitate at any given temperature.
Austenite grain growth behaviour in the steels after deformation
Figure 1 – Schematic representation of the deformation process on the Bähr dilatometer
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Figure 4 shows the grain size distribution in the Nb-bearing steels after high-temperature deformation followed by tempering at 490°C for 72 hours. The steels showed a unimodal grain structure during deformation. These distributions and structures attest to the fact that no abnormal grain growth, i.e. no bimodal distribution, was visible in these steels at the thermomechanical controlled processing (TMCP) conditions employed in this work, confirming the presence of pinning particles in the steels as predicted by Thermo-CalcTM. The Journal of The Southern African Institute of Mining and Metallurgy
The influence of niobium content on austenite grain growth in microalloyed steels
Figure 2 – Thermo-CalcTM predictions of the volume fraction of precipitates in the high-N laboratory-produced steels with (a) 0.03 wt%, (b) 0.05 wt%, (c) 0.07 wt%, and (d) 0.1 wt% Nb
basic relationship for the pinning force of particles on grain boundaries
Increased Nb content and austenite grain growth in the microalloyed steels The observed effect of an increased Nb content in the microalloyed steels is shown in Figure 5. It is evident that there was limited grain growth in the steel containing 0.1 wt% Nb, as it produced the smallest final grain sizes, while the highest grain growth was seen in the 0.002 wt% Nb steel, thereby following the volume fraction of Nb precipitates in the steel. The observed results are in agreement with Zener’s The Journal of The Southern African Institute of Mining and Metallurgy
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Figure 3 – Effect of Nb content on volume fraction of niobium precipitate as predicted by Thermo-calcTM
(Zener, 1948, as cited by Ringer, Li, and Easterling, 1989), which predicts a stronger effect of pinning with increased volume fractions of small precipitates at constant particle sizes. The higher Nb content accounts for a greater volume fraction of NbN precipitates as shown in Figure 3, providing a greater area fraction of solute-rich regions compared to the steels containing less Nb (Xu and Thomas, 2011; Brewer, Erven and Krauss, 1991; Deus et al., 2002; Akamatsu, Senuma and Hasebe, 1992). The quantitative analysis of the optical measurements of austenite grain growth, is shown in Figure 6, which confirms the substantial influence of an increased Nb content on austenite grain growth inhibition in these steels. Quantitative evaluations of the average grain size as a function of austenitizing time and temperature in the 0.03 wt% Nb microalloyed steel are shown in Figures 7 and 8 respectively.A comparative analysis of Figures 7 and 8 indicates the expected greater effect of temperature compared with time on grain growth, as recorded by numerous studies (Seok et al., 2014; Yue et al., 2010; Nanba et al., 1992; Akamatsu, Senuma and Hasebe, 1992; Sha and Sun, 2009; Zhao et al., 2011)
The influence of niobium content on austenite grain growth in microalloyed steels
Figure 4 – Optical micrographs and grain size distribution for (a) an 0.07 wt% Nb deformed at 1100°C at a strain rate of 0.1 s-1 to a strain of 0.4, (b) 0.05 wt% Nb deformed at 1150°C at a strain rate of 0.1 s-1 to a strain of 0.4
Figure 5 – Optical micrographs of (a) 0.002 wt%, (b) 0.03 wt%, (c) 0.05 wt%, (d) 0.07 wt%, and (e) 0.1 wt% Nb steels deformed at a temperature of 1100°C and held for 120 minutes
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The influence of niobium content on austenite grain growth in microalloyed steels Development of a constitutive equation It is generally accepted that the grain boundary migration velocity (v) is proportional to the driving force (Seok et al., 2014; Maalekian et al., 2012; Yue et al., 2010; Olasolo et al., 2011; Alogab et al., 2007; Hodgson and Gibbs, 1992; Nanba et al., 1992), an approximation justified at the relatively small driving forces involved in grain growth. This velocity is given by the expression (Stumpf, 2010) [1] where M is the grain boundary mobility, Fd is the driving force for grain growth, and Fp is the pinning force induced by precipitates. From the rate of movement: Figure 6 – Influence of Nb content on the average austenite grain size (AGS), showing a strong inhibition in Nb-containing steels
[2] The grain boundary mobility M is given by the expression
and the grain growth rate equation therefore becomes: [3] Under isothermal conditions [4] The differential solution to Equation [4] is given by (Sellars and Whiteman, 1979) [5] Figure 7 – Isothermal grain growth behaviour of austenite in steels, showing an increase in the AGS with time in the 0.03 wt% Nb steel after hot deformation
where T is the absolute temperature, t is the time, Q is the activation energy for grain boundary migration, A is a constant parameter dependent on the material and processing conditions, and R is the universal gas constant. A general solution to Equation [4] in a linear form is expressed in Equation [6]: [6] Equation [5] has been used by many authors (Seok et al., 2014; Maalekian et al., 2012; Nanba et al., 1992; Florez and Martinez, 1997; Rollett, Srolovitz and Anderson, 1989; Sellars and Whiteman, 1979; Fu et al., 2011; Wang and Wang, 2008; Wang et al., 2006) by assuming the value of Do to be zero or constant. This simplifies the plot of lnD as a function of 1/T. In the current study Do, which was experimentally measured, was not discarded or kept constant but was included and used in Matlab programming and an Excel solver to plot
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Figure 8 – Isothermal grain growth behaviour of austenite in steels, showing an increase in the AGS with temperature in the 0.03 wt% Nb steel after initial hot deformation at the same temperature
as a function of 1/T in generating the constants used in the development of the constitutive equations. This was done through regression analyses of the first and second derivative orders of the experimentally derived regression value R2.
The influence of niobium content on austenite grain growth in microalloyed steels The grain growth constants n, Q, and A The constant n was obtained by regression analysis of experimentally measured grain sizes through iterative tests of the straight-line regression coefficients R2, from a plot of Equation [6] where a second differential of the R2, (dR2)2/dn2 with respect to n gave a maximum correlation value of n at (dR2)2/dn2 = 0. Figure 9 shows a plot of ln(−) as a function of the inverse temperature 1/T, which was used in determining the constants. Thus, a plot of Equation [6] provided a linear relationship between
and 1/T. Comparison of Equation [6] with the equation of a straight line made it possible for the constants Q and A to be determined from the plots shown in Figure 9. Thus,
Figure 9 – Plot of the natural log of the average grain size as a function of the inverse of the absolute temperature for the 0.07 wt% Nb steel
(while the slope is negative) and A = ec/t where, c is the intercept from the graph and t is the time. The dependence of the grain growth constants n, Q, and A on the Nb content was quantitatively analysed by multiple regression to obtain the following reliance equations, where [Nb] is the wt% in the steel:
growth equation for austenite grain growth prediction in 0.002–0.1% Nb microalloyed steels during deformation within the temperature range of 1000–1150°C and delay times ranging from 0–120 minutes is constituted for DDef in μm:
The statistical analysis resulting from the multiple regression analysis presented in Table II, showing that the activation energy Q has a strong linear relationship with the Nb content, while there is no linear relationship between A and the Nb content. There is, however, a weak linear relationship between the grain growth constant n and the Nb content. The constants derived from experimentally measured grain sizes for the microalloyed steels under deformation conditions with and without taking Do into consideration are presented in Table III.
where Do is the initial austenite grain size in μm and [Nb] is in wt %.
Table II
Statistical results of the multiple regression analysis of the dependence of the austenite grain growth constants on varied Nb content Regression statistics
General constitutive equation for austenite grain growth prediction in Nb-bearing microalloyed steels with high N contents
Adjusted R square Standard error Number of observations ANOVA
Based on the analysis of the isothermal grain growth kinetics in Nb-microalloyed steels, the following isothermal grain
Grain growth constants Q 0.91 32.31 5 0.0072***
n 0.64 0.83 5 0.066*
A 0.40 1.7 × 1028 5 0.152
*** p<0.01 (this means there is a very strong relationship with probability <1%) * p<0.1 (this indicates a weak relationship with probability <10%)
Table III
Grain growth constants generated from deformation data of Nb-microalloyed steels Constants determined from deformed Nb-microalloyed steels Steel (wt % Nb)
0.002 0.03 0.05 0.07 0.1
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Constants using measured values of Do
Constants with Do= 0
Difference in constants from neglecting D0
n
A
Q
n
A
Q
Δn
ΔA
ΔQ
2.8 5.5 5.6 6.0 6.2
4.24E+12 1.33E+22 7.94E+23 9.63E+24 4.96E+28
276 417 474 489 572
2.8 5.5 5.5 6.1 6.2
3.00E+11 6.60E+21 5.00E+23 4.62E+24 1.22E+28
239 384 428 459 559
0 0 0.1 -0.1 0
3.94E+12 6.70E+21 2.94E+23 5.01E+24 3.74E+28
37.27 33.10 46.31 30.41 12.67
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The influence of niobium content on austenite grain growth in microalloyed steels Discussion Austenite grain growth behaviour of Nb-bearing microalloyed steels The grain growth rate was significantly slower in microalloyed steels with higher contents of Nb. Grain growth control in microalloyed steels has been shown to be dependent on the amount of microalloying additions (Alogan et al., 2007; Hodgson and Gibbs, 1992). The greatest inhibition here was recorded in the steel containing 0.1 wt% Nb. It was also found that the average austenite grain size decreased with increasing Nb content in the steels. Generally, it has been reported that increasing the N content in the steel leads to a beneficial increase in grain refinement (Alogan et al., 2007; Hodgson and Gibbs, 1992; Nanba et al., 1992), which without doubt results from the larger volume fraction of the precipitates that act as pinning particles. It should be noted, however, that while increased N content is beneficial for grain refinement, excessive levels of N dissolved in austenite may be detrimental to other properties of the steel, such as hardenability, through the decrease of the Nb fraction dissolved in the austenite, which is a key in improving most of the desired mechanical properties of the steel (Alogan et al., 2007; Hodgson and Gibbs, 1992; Nanba et al., 1992).
Predictive potential of the equation for grain growth in microalloyed steels in the current study A logical degree of precision in predicting austenite grain growth in Nb bearing steels has been achieved by comparison of experimentally measured grain sizes with predicted grain sizes using the general constitutive equations developed in
this study, as shown in Figure 10. The correlation R2 values obtained from the comparison of the experimentally measured grain sizes with the predicted grain sizes are 0.94, 0.91, and 0.90 for 0.03 wt% Nb, 0.05 wt% Nb, and 0.07 wt% Nb steels respectively. The 0.1 wt% Nb steel, however, recorded a poor correlation value of R2 = 0.63, showing the inability of the equation to predict grain sizes for steels containing 0.1 wt% Nb and also most likely for those with higher Nb contents. This is due to the inability of the developed equation to account for the excess precipitates that form owing to the high addition of the microalloying element.
Conclusions ➤ A general constitutive equation for predicting austenite grain growth in Nb-bearing steels that incorporates the Nb content and then the starting grain size Do after simulated rough rolling has been developed ➤ There is a linear relationship between Nb content and the activation energy for grain growth ➤ The austenite grain growth exponent n, the activation energy Q, and the material and processing conditions constant A all increase with an increase in Nb content, most likely from increased effectiveness of grainboundary pinning in the microalloyed steel with increased Nb content ➤ The constitutive equation shows a poor correlation for the steel containing 0.10 % Nb, with an R2 value of 0.63 recorded in the plot of measured against predicted grain sizes for the 0.1 wt% Nb steel. The equation is most likely not applicable at such high values of Nb.
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Figure 10 – Comparison of predicted and measured austenite grain sizes in Nb-bearing steels
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http://dx.doi.org/10.17159/2411-9717/2015/v115n10a13
The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel by H.J. Uananisa*, C.W. Siyasiya*, W.E. Stumpf* and M.J. Papo†
The need to reduce weight while maintaining good mechanical properties in materials used in the automotive industry has over the years seen an increased exploitation of various steels to meet this new demand. In line with this development, the ferritic stainless steel family has seen a wide application in this industry, with the AISI 436 type increasingly being used for automotive trims and mufflers for exhaust systems, as well as a significant part of this steel’s application being for the manufacture of wheel nuts and wheel nut caps in trucks, mainly through the deep drawing process. However, there have been reports of some poor surface roughening of this material during deep drawing, with tearing and/or cracking also reported in some instances. This has been suspected to possibly be associated with some local differences in localized mechanical properties between grains and grain clusters of the rolled and annealed material. In order to investigate the poor surface roughness exhibited by AISI 436 ferritic stainless steel (FSS) during deep-drawing, Lankford values (Rmean and Δr), grain size, and microtextures of various sheet samples from this steel were studied. The chemical composition range for the samples was 0.013–0.017% C, 17–17.4% Cr, 0.9–1% Mo, and 0.4–0.5% Nb. The steels were subjected to various hot and cold rolling processing routes i.e. involving industrial direct rolling (DR) or intermediate annealing rolling (IR), and the drawability and final surface qualities of the steels were compared. It was found that the DR route gave an average R-mean and Δr value of 1.9 and -1.4 respectively, while the IR route yielded an average Rmean and Δr value of 1.6 and 0.52 respectively. The high Δr value for the DR processing route had a substantial adverse effect on the drawability. IR samples exhibited a smoother surface finish on visual inspection, while clear flow lines were visible on the DR samples, despite the fact that DR is the preferred industrial processing route due to the reduced production costs it offers. This observation was also confirmed through SEM examinations. The difference in the surface quality was attributed to microtexture. However, the mechanism responsible for this difference still needs to be identified. Keywords ferritic stainless steel, formability, microtexture, EBSD.
Introduction Owing to the price volatility of nickel over recent years, the need to cut down on the use of nickel in the steelmaking industry has become a matter of keen interest, hence the rising interest in ferritic stainless steels to replace their austenitic counterparts for various industrial applications. Ferritic grades, containing chromium and possibly other stabilizing elements (Ti, Mo, Nb, etc.) are well known as cost-saving materials as they do not have the expensive nickel additions (Charles et The Journal of The Southern African Institute of Mining and Metallurgy
al., 2008), while at the same time presenting good mechanical properties similar to those offered by austenitic steel grades. Moreover, some standard ferritic grades such as 409, 410, and 430 are readily available all over the world, and are already very successfully used for various applications, such as washingmachine drums and exhaust systems in the automotive industry, and actually have a much broader application potential in many fields (Charles et al., 2008). Since new designs in the automotive industry are driven by safety and environmental concerns, the need to reduce the weight of the exhaust systems while maintaining good resistance to thermal fatigue is continuing to see the exploitation of the ferritic stainless steel family for alternative solutions. Ferritic stainless steels have a relatively low coefficient of thermal expansion and, therefore, some efforts have been made to create new ferritic stainless steels with high yield strengths at elevated temperatures, particularly by the addition of niobium (Nb), which increases the initial high-temperature strength through solid solution hardening (Sello and Stumpf, 2010). The ferritic type AISI 436 is increasingly used for automotive trims, with a major application being for lug nuts or wheel nuts on trucks. It has also seen extensive use in the production of both the central and rear mufflers of exhaust systems, due to its good hot and wet corrosion resistance properties (Charles et al., 2008). However, during the deep drawing process associated with the manufacture of most of these products, tearing or cracking can sometimes occur. Deep drawing is one of the most common processes for forming metal parts from sheet metal plates, and it is widely used for the mass production of parts in the
* University of Pretoria, Department of Material Science & Metallurgical Engineering, South Africa. † Advanced Materials Division, Mintek. © The Southern African Institute of Mining and Metallurgy, 2015. ISSN 2225-6253. Paper received Aug. 2015. VOLUME 115
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Synopsis
The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel emphasis on the processing route, influence this poor surface quality behaviour of AISI 436. This was achieved by studying some crystallographic parameters of the steel, and correlating those parameters to surface topography, in order to help understand the mechanism behind the surface roughening of the steel in subsequent studies. The ultimate objective is to develop possible solutions to the problems experienced in the deep drawing of AISI 436. Figure 1 – A multi-stage deep-drawing process for wheel-nut covers
Experimental procedure
packaging industry, automotive industry, and many others. A typical example of this process is illustrated in Figure 1, which shows some wheel-nut covers manufactured through a six-stage deep-drawing process. In the reported cases of poor surface roughness, tearing, and cracking, the surface of the drawn products appears to have roughened, possibly due to local differences in mechanical properties between grains or grain clusters of the raw material. This behaviour has been suspected to be a crystallographic texture effect similar to ridging and roping, which is the very common susceptibility of ferritic stainless steel to develop narrow ridges on the sheet surface during forming (Knutsen and Wittridge, 2002; Raabe et al., 2003; Shin et al., 2003). The ridges or undulations result in a very dull surface appearance, which in turn reduces the surface shine and quality of the formed product (Knutsen and Wittridge, 2002). The purpose of this study is to investigate how various thermomechanical processes, with particular
Various samples of commercial AISI 436 were received and characterized. These were sampled from a production trial operation, as shown in Figure 2, in which two processing routes, direct rolling (DR) and intermediate rolling (IR), were investigated. The emphasis in this study was limited to the final cold-rolled and annealed samples marked F1, F2, F3, and F4 in Figure 2, with thicknesses ranging from 0.46 mm to 0.50 mm. All the samples were processed from the same ‘production heat’ and hence had the same chemical composition, shown in Table I. Each sample was first mechanically polished, etched in an aqua regia solution, and examined under an optical microscope to determine the grain structure morphology. Grain-size measurements were determined by the ‘Image J’ line-intercept method. The samples were then further mechanically prepared to a final fine polish in a colloidal silica medium (OP-S) for electron backscatter diffraction (EBSD) analysis, with the scans being performed using a Jeol JSM-IT300LV scanning electron microscope (SEM).
Figure 2 – Block diagram showing the thermomechanical trial processing routes for the AISI 436 samples used in this work
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The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel Table I
Chemical composition of the AISI 436, wt.% C
Mn
Si
P
S
Cr
Mo
Ni
Al
Cu
Nb
Ti
Sn
N
0.013
0.42
0.43
0.022
0.002
17.14
0.94
0.25
0.003
0.09
0.41
0.001
0.012
0.017
Table II
Chemical composition of the deep-drawn AISI 436 samples, wt.% Sample
C
Mn
Si
P
S
Cr
Mo
Ni
V
Cu
Nb
Ti
Co
N
D1 D2
0.015 0.017
0.46 0.56
0.55 0.38
0.02 0.02
0.0005 0.0005
17.04 17.23
0.94 0.90
0.15 0.24
0.12 0.10
0.08 0.10
0.50 0.37
0.002 0.003
0.02 0.02
0.02 0.02
Orientation distribution functions (ODFs) were further calculated by the series expansion method (lmax=22) using the SALSA post-processing software of the CHANNEL5 package. The mean R-values (Rm) and planar anisotropy values (Δr) for the various samples were measured after 10% strain along the longitudinal (0°), transverse (90°), and diagonal (45°) directions. The Rm and Δr values were then calculated using the following standard equations:
[1]
of the deep-drawn items (Maruma et al., 2013). Samples of the steel from different production heats, with the compositions shown in Table II, were then deep-drawn into lug-nut covers through the multi-stage forming process shown in Figure 1, and analysed for surface roughness. Sample D1 was intermediate rolled (IR) prior to the deepdrawing process, while D2 was directly rolled (DR). This analysis was used to investigate the effect of the production processing route on the surface roughness of the steel.
Results and discussion Microstructural analysis
[2] The subscripts 0°, 45°, and 90° refer to the longitudinal, diagonal, and transverse directions with respect to the rolling direction. The planar anisotropy (Δr) gives an indication of the amount of necking or earing that will occur on the edges
Figure 3 shows the microstructures of the cold-rolled and annealed samples by optical microscopy. Grain size measurements showed an average grain size value of about 25 μm for the four final cold-rolled and annealed samples (F1 to F4), with the respective measured values shown in Table II. It is evident from these results that the processing route
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Figure 3 – Microstructures of the samples after a 30-second deep etching in aqua regia solution. (a) F1 sample, (b) F2 sample, (c) F3 sample, and (d) F4 sample
The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel Table III
R-mean, Δr, and measured grain sizes Sample
RHF Dis. temp (°C)
Rolling route
Annealing after Steckel
Rm
delta-r
F1 F2 F3 F4
1079 1083 1079 1160
IR IR DR DR
No Yes No No
1.55 1.55 1.65 2.07
0.53 0.51 -1.38 -1.39
had no significant effect on grain size as the average grain sizes measured were within an average ±3 μm standard deviation from each other. The optical microstructures also clearly reveal various grains that are strongly etched (darker) for all samples, while others are lightly etched. This disparity in grain morphology is caused by differences in grain orientations (Raabe and Lucke, 1992). In order to evaluate the drawability of the sheet samples, measurements of the R-mean (Lankford parameters – Rm) values for each sheet were determined through tensile tests. The results, also shown in Table III, indicated a lower average Rm value for the IR route at about 1.6 compared to an average Rm value of 1.9 for the DR route. It is of particular interest that the planar anisotropy parameter (Δr) for both DR samples is significantly higher in magnitude and negative (average value of -1.4) compared to that of the IR samples (0.52). This behaviour is indicative of a rotation of earing from an angle of 0° and 90° with respect to the rolling direction (RD) to a 45° or 135° direction with respect to RD, which could be attributed to the poor resulting formability of the DR samples despite their higher Rm value. The severity of the surface roughness of the DR samples was further illustrated by examining one sample from each route by SEM. The micrographs are shown in Figure 4. Both samples were obtained from the ‘wall area’ of a stage two sample (as shown in Figure 1) from each route, with D1 being from the IR process and D2 from the DR process. It is evident, even by visual inspection of this early forming stage of the two samples, with scans at the same magnification, that the DR sample shows severe surface roughness compared to the flatter surface for the IR sample. Moreover,
Grain size (μm) 24.0 ± 3.7 27.6 ± 3.8 24.2 ± 2.1 22.0 ± 2.5
as may be seen from the distribution of the ‘hills and valleys’ on the micrographs, some areas on the IR sample seem not to have undergone any deformation (circled regions on ‘D1’ in Figure 4). This observation calls attention to a phenomenon mentioned by Knutsen and Wittridge (2002), who emphasize the influence of grain size banding in the light of the HallPetch effect, hence possible resultant differential yielding under tension. That is to say, clusters consisting of different grain sizes would obviously respond to deformation differently, thereby resulting in distortion of the surface.
Texture evolution Many studies have shown that ferritic stainless steels, which have a body-centred cubic (bcc) structure, have a tendency to form a preferred orientation of their grains (fibre texture) during rolling and annealing. Siqueira et al. (2011), Huh et al., (2005), and Maruma Siyasiya, and Stumpf (2013) have elaborated on the fact that rolling deformation in most cases leads to a texture characterized by two orientation fibres – an α-fibre texture typically comprising orientations with a common <110> direction parallel to RD (RD//<110>, and a γfibre texture comprising orientations of the {111} plane parallel to the RD. Normally, subsequent annealing of coldrolled sheets increases the γ-fibre component at the expense of the α-fibre, with possible improvements in the formability of the sheet metal (Yazawa et al., 2003). Figure 5 shows the orientation distribution function (ODF) maps of the four samples in this work alongside a Bunge notation texture diagram showing the main texture fibres in bcc taken at Φ=45°. The F1 and F2 IR sample textures are predominantly characterized by a strong γ-fibre, notably {111}<011>, {111}<123>, and {111}<112>, as well as {554}<225>. The
Figure 4 – SEM micrographs of deep-drawn AISI 436 sheets after the second stage of the wheel-nut cover forming process. D1: smoother surface of IR sample, and D2: severe surface roughness of DR sample
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The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel striking difference between the two samples is observed in the intensities of these γ-fibres, with the F1 sample showing a stronger intensity of a maximum of about 11 compared to a maximum of about 5 for the F2 sample. This observation clearly illustrates the insignificance of the AP1 annealing process, which is clearly an additional industry cost. The F3 DR sample, however, is characterized by a strong intensity near or between {322}<258> and {322}<236>. Thus, it is evident that a strong texture component in the {100} plane would always have an adverse influence on the
annealing texture (Maruma et al., 2013). This is possibly compounded by the absence of subsequent ‘intermediate annealing’ in the DR process, which hindered an increase of the γ-fibre evolution. Similarly, the F4 DR sample shows a weak α-fibre near {322}<258> as well as a weak γ-fibre in the vicinity of {554}<225>, despite it having the highest Lankford parameter value (Rm = 2.1). The high negative earing parameter (Δr), which has already been alluded to, is suspected of adversely affecting the formability of these sheets (F3 and F4), despite their good ductility parameters.
The Journal of The Southern African Institute of Mining and Metallurgy
VOLUME 115
OCTOBER 2015
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Figure 5 – Orientation distribution function (ODF) maps for the four samples studied. The adjacent diagram shows the texture fibre positions in Euler space (Raabe and Lucke, 1992)
The influence of thermomechanical processing on the surface quality of an AISI 436 ferritic stainless steel
Figure 6 – Inverse pole figure maps //ND for the first stages of the deep-drawn samples, showing clear indications of grain size clustering
Effect of thermomechanical processing on grain size distribution The grain size clustering effect mentioned above is clearly illustrated by the inverse pole figure (IPF) maps in Figure 6, particularly on the map for sample D2 (the DR sample), in which there is evident clustering of smaller grains separately in a particular region, almost forming a band of smaller grains, and clustering of larger grains in other regions within the microstructures. However, despite this grain size clustering, which may be responsible for differences in yielding across the sheet metal plates and hence the poor surface roughness, no clear texture banding along a particular direction was observed in the microstructures or the SEM images, as is normally the case in ridging and roping. This leads to the conclusion that another localized grain-related mechanism, rather than ridging, is at play in this case.
CHARLES, J., MITHIEUX, J.D., SANTACREU, P.O. and PEGUET, L. 2008. The ferritic stainless steel family: the appropriate answer to nickel volatility. 6th European Stainless Steel Conference, Helsinki, 10–13 June 2008. G. SEARCH, “Euler space - recrystallization textures of metals,” [Online]. Available: www.google.com. [Accessed 12 May 2015]. HUH, M.-Y., LEE, J.-H., PARK, S.H., ENGLER, O. and RAABE, D. 2005. Effect of through-thickness macro and micro-texture gradients on ridging of 17% Cr ferritic stainless steel sheet. Materials Technology - Stainless Steels, vol. 11, no. 76. pp. 797–806. KNUTSEN, R.D. and WITTRIDGE, N.J. 2002. Modelling surface ridging in ferritic stainless steel. Materials Science and Technology, vol. 18. pp. 1279–1285. MARUMA, M.G., SIYASIYA, C.W. and STUMPF, W.E. 2013. Effect of cold reduction and annealing temperature on texture evolution of AISI 441 ferritic stainless steel. Journal of the Southern African Institute of Mining and
Conclusions
Metallurgy, vol. 113, no. 2. pp. 115–120.
➤ Grain size measurements suggested that the processing route had no significant influence on average grain size, but had an effect on grain size distribution ➤ The intermediate rolling (IR) route resulted in the desired γ-fibre texture, and hence superior surface qualities (surface roughness) in comparison to the direct rolling (DR) route ➤ In addition, the IR route also resulted in lower planar anisotropy (average Δr = 0.52) compared to the DR route (Δr = -1.4), suggesting superior deep-drawability as observed ➤ Grain-size clustering as opposed to texture banding is suspected as a possible factor responsible for surface roughening.
Acknowledgements
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OCTOBER 2015
RAABE, D. and LUCKE, K. 1992. Influence of particles on recrystallization textures of ferritic stainless steels. Steel Research, vol. 63, no. 10. pp. 457–462. RAABE, D., SACHTLEBER, M., WEILAND, H., SCHEELE, G. and ZHAO, Z., 2003. Grainscale micromechanics of polycrystal surfaces during plastic straining. Acta Materialia, vol. 51. pp. 1539–1560. SELLO, M.P. and STUMPF, W.E. 2010. Laves phase embrittlement of the ferritic stainless steel type AISI 441. Materials Science and Engineering Section A, vol. 527. pp. 5194–5202. SHIN, H.-J., AN, J.-K., PARK, S.H. and LEE, D.N. 2003. The effect of texture on ridging of ferritic stainless steel. Acta Materialia, vol. 51. pp. 4693–4706. SIQUEIRA, R.P., SANDIM, H.R.Z., OLIVEIRA, T.R. and RAABE, D. 2011. Composition
The authors gratefully acknowledge the financial contribution provided by the Advanced Metal Initiative (AMI) of the Department of Science and Technology (DST) through the Ferrous Metals Development Network (FMDN), and Columbus Stainless (Middelburg, South Africa). Special thanks are also due to Mr Dave Smith and Mr Jaco Kruger (Columbus Stainless) for their immense contribution and input into this study. 986
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and orientation effects on the final recrystallization texture of coarsegrained Nb-containing AISI 430 ferritic stainless steels. Materials Science and Engineering Section A, vol. 528, no. 9. pp. 3513–3519. YAZAWA, Y., OZAKI, Y., KATO, Y. and OSAMU, F. 2003. Development of ferritic stainless steel sheets with excellent deep drawability by {111} recrystallization texture control. Society of Automotive Engineers (SAE) of Japan, vol. 24. pp. 483–488.
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The Journal of The Southern African Institute of Mining and Metallurgy
14â&#x20AC;&#x201C;17 March 2016, Gaborone International Convention Centre Workshop: 14 March 2016 Conference: 15â&#x20AC;&#x201C;16 March 2016 Technical Visit: 17 March 1026 BACKGROUND eing the sixth conference in the series, the Diamondsâ&#x20AC;&#x201D;still Sparkling Conference targets the full spectrum of the diamond pipeline from exploration through to sales and marketing. The last conference was held in 2013 at Misty Hills, Muldersdrift: the 2016 conference is returning to Botswana which previously hosted in 2010.
Photographs of diamonds courtesy of Petra Diamonds
B
he objective of the conference will be to provide a forum for the dissemination of information relating to the latest mining methods and technologies applicable to the diamond mining industry. This will consider all stages of the value chain, from exploration through mine design, drilling and blasting production, and processing, to cutting, marketing and sales.
T
> Processing engineers > Mining engineers > Geotechnical engineers > Geologists > Consultants > Suppliers
> > > > > >
Sponsors:
Cutting and polishing Marketing and sales Diamontiers Mine managers Mining companies Students mining industry
> Geology and exploration > Mine expansion projects > Mining, metallurgical and beneficiation technology > Rough diamond sales and marketing > Cutting and polishing > Financial services and industry analysis > Industry governance and legislation update > Mine specific case studies
Conference Supporter Media Partner
For further information contact: Yolanda Ramokgadi â&#x20AC;˘ Conferencing co-ordinator ¡ SAIMM, Pâ&#x20AC;&#x2C6;Oâ&#x20AC;&#x2C6;Box 61127, Marshalltown 2107 Tel: (011) 834-1273/7 â&#x20AC;˘ Fax: (011) 833-8156 or (011) 838-5923 E-mail: yolanda@saimm.co.za â&#x20AC;˘ Website: http://www.saimm.co.za
INTERNATIONAL CHAIR Dr. Raj K. Singhal (singhal@shaw.ca) HONORARY CHAIR Mr. Mike Teke CHAIRMAN MPES 2015 Prof. Cuthbert Musingwini
3 ECSA POINTS ESSES: KEYNOTE ADDR
M. Teke N. Froneman G. Lane man R. Webber-Young
LOCAL ORGANISING COMMITTEE Dr Bekir Genc Dr Steven Rupprecht Mr Alastair Macfarlane Mr Kelello Chabedi Mr Jannie Maritz Mr Godknows Njowa Mr Mike Woodhall Prof. Jim Porter Mr Alex Bals Dr Andre Dougall Dr Gordon L. Smith Dr Craig Smith Mr I. Wermuth Mr C. Birch Mr G. Lane CO-CHAIRS Professor Monika Hardygóra Prof. Carsten Drebenstedt Prof. Uday Kumar Prof. Kostas Fytas Prof. Petr.Sklenicka INTERNATIONAL ORGANISING COMMITTEE Prof. Hani Mitri Dr Nuray Demirel Dr Marilena Cardu Dr Fiona Mavroudis Dr Meimei Zhang Prof. Ge Hao Prof. Celal Karpuz Prof. Liu Mingju Dr Mohan Yellishetty Prof. Hideki Shimada Dr Gento Mogi Dr Vera Muzgina Dr Morteza Osanloo Dr Juri- Rivaldo Pastarus Prof. Hakan Schunnesson Ms. M. Singhal Dr Eleonora. Widzyk-Capehart Prof. Antonio Nieto Prof. Michael A. Zhuravkov Prof. Sukumar Bandopadhaya Prof. Ferri Hassani Dr Noune.Melkoumian Dr Joerg Benndorf Prof. Giorgio Massacci Dr Maria Menegaki Prof. Svetlana V.Yeffremova Prof. Pivnyak Gennadiy Prof. Vladimir Kebo Dr Marie Vrbova
23rd INTERNATIONAL SYMPOSIUM on MINE PLANNING & EQUIPMENT SELECTION
MPES 2015
Smart Innovation in Mining 8 November 2015—Cocktail Function 9–11 November 2015—Conference 12 November 2015—Tours and Technical Visits Sandton Convention Centre, Johannesburg, South Africa BACKGROUND The Southern African Institute of Mining and Metallurgy (SAIMM) will be hosting the 23rd International Symposium on Mine Planning and Equipment Selection (MPES) for the first time in South Africa in 2015. The symposium has successfully been organized annually for the past 25 years with intermittent 1-year recess as necessary. In 2013 the symposium was held in Germany. It will be in recess in 2014. MPES has previously been held in Turkey, Greece, Canada, Kazakhstan, Australia, Czech Republic, Brazil, India, China, Ukraine, Poland and Italy. Other venues for future MPES are Czech Republic, Sweden, and Australia.
OBJECTIVES The key objective of MPES is to provide a platform for researchers from academic institutions, professionals from mining companies, practitioners from consulting companies, equipment suppliers (OEMs) and software providers to share the latest global developments in mine planning and equipment selection across all commodities for the benefit of the mining industry in improving efficiencies and safety. An additional objective of MPES 2015 is to encourage MSc and PhD students to showcase their research in a ‘Young Authors Category’ to foster creation of the next generation of MPES participants.
MAJOR THEMES TO BE COVERED ALONG THE VALUE CHAIN • • • • • • • • • • • • • • •
Data Collection and Modelling: State of the Art Practices Mineral Resource and Mineral Reserve Estimation and reporting Economic and Technical Feasibility Studies, Mine Development Case Studies Design, Planning and Optimization of Surface and Underground Mines Transition from surface to underground mining Rock Mechanics and Geotechnical Applications Mining Equipment: Selection, Operation, Control, Monitoring and Optimization Mechanization and Automation of Mining Processes Application of Information Technology Short interval/planning and control Resource-to-Market: Reconciliation and Optimization Productivity and Competitiveness of Mining Operations Sustainability: Improving Health, Safety and Environmental Practice and Performance Mine Closure and Rehabilitation in mine planning Young Authors Category (MSc/PhD Students below 35 years).
Confirmed Technical Visits: Please note that a maximum of 20 delegates can be accommodated on each of these Technical Visits • Zibulo Colliery (Underground Coal)
• Grootegeluk Mine (Open Pit Multi-Seam)
• Bathopele Mine (Mechanised Platinum)
• Tau Tona Mine (Deep Level Gold Mine)
• University of Pretoria Virtual Reality Laboratory
Sponsor:
EXHIBITION/SPONSORSHIP Sponsorship opportunities are available. Companies wishing to sponsor or exhibit should contact the Conference co-ordinator.
SHELL PECTEN
CONFERENCE ANNOUNCEMENT
For further information contact, please contact: Dr Raj Singhal, E-mail: singhal@shaw.ca or Conference Organisers: The Southern African Institute of Mining and Metallurgy (SAIMM) · E-mail: raymond@saimm.co.za · Official Website: http://www.saimm.co.za
INTERNATIONAL ACTIVITIES 2015
20 October 2015 â&#x20AC;&#x201D; 13th Annual Southern African Student Colloquium Mintek, Randburg, Johannesburg Contact: Yolanda Ramokgadi Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: yolanda@saimm.co.za Website: http://www.saimm.co.za 21â&#x20AC;&#x201C;22 October 2015 â&#x20AC;&#x201D; Young Professionals 2015 Conference Making your own way in the minerals industry Mintek, Randburg, Johannesburg 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 28â&#x20AC;&#x201C;30 October 2015 â&#x20AC;&#x201D; AMI: Nuclear Materials Development Network Conference Nelson Mandela Metropolitan University, North Campus Conference Centre, Port Elizabeth Contact: Raymond van der Berg Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: raymond@saimm.co.za Website: http://www.saimm.co.za 8â&#x20AC;&#x201C;12 November 2015 â&#x20AC;&#x201D; MPES 2015: Twenty Third International Symposium on Mine Planning & Equipment Selection Sandton Convention Centre, Johannesburg, South Africa Contact: Raj Singhal E-mail: singhal@shaw.ca or E-mail: raymond@saimm.co.za Website: http://www.saimm.co.za
2016 14â&#x20AC;&#x201C;17 March 2016 â&#x20AC;&#x201D; Diamonds still Sparkle 2016 Conference Gaborone International Convention Centre Contact: Yolanda Ramokgadi Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: yolanda@saimm.co.za Website: http://www.saimm.co.za
13â&#x20AC;&#x201C;14 April 2016 â&#x20AC;&#x201D; Mine to Market Conference 2016 South Africa Contact: Yolanda Ramokgadi Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: yolanda@saimm.co.za Website: http://www.saimm.co.za 17â&#x20AC;&#x201C;18 May 2016 â&#x20AC;&#x201D; The SAMREC/SAMVAL Companion Volume Conference Johannesburg Contact: Raymond van der Berg Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: raymond@saimm.co.za Website: http://www.saimm.co.za 21â&#x20AC;&#x201C;28 May 2016 â&#x20AC;&#x201D; ALTA 2016 Perth, Western Australia Contact: Allison Taylor Tel: +61 (0) 411 692 442 E-mail: allisontaylor@altamet.com.au Website: http://www.altamet.com.au May 2016 â&#x20AC;&#x201D; PASTE 2016 International Seminar on Paste and Thickened Tailings Kwa-Zulu Natal, South Africa Contact: Raymond van der Berg Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: raymond@saimm.co.za Website: http://www.saimm.co.za 9â&#x20AC;&#x201C;10 June 2016 â&#x20AC;&#x201D; 1st International Conference on Solids Handling and Processing A Mineral Processing Perspective South Africa Contact: Raymond van der Berg Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: raymond@saimm.co.za Website: http://www.saimm.co.za 1â&#x20AC;&#x201C;3 August 2016 â&#x20AC;&#x201D; Hydrometallurgy Conference 2016 â&#x20AC;&#x2DC;Sustainability and the Environmentâ&#x20AC;&#x2122; in collaboration with MinProc and the Western Cape Branch Cape Town Contact: Yolanda Ramokgadi Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: yolanda@saimm.co.za Website: http://www.saimm.co.za 16â&#x20AC;&#x201C;19 August 2016 â&#x20AC;&#x201D; The Tenth International Heavy Minerals Conference â&#x20AC;&#x2DC;Expanding the horizonâ&#x20AC;&#x2122; Sun City, 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
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12â&#x20AC;&#x201C;14 October 2015 â&#x20AC;&#x201D; Slope Stability 2015: International Symposium on slope stability in open pit mining and civil engineering In association with the Surface Blasting School 15â&#x20AC;&#x201C;16 October 2015 Cape Town Convention Centre, Cape Town Contact: Raymond van der Berg Tel: +27 11 834-1273/7, Fax: +27 11 838-5923/833-8156 E-mail: raymond@saimm.co.za Website: http://www.saimm.co.za
Company Affiliates The following organizations have been admitted to the Institute as Company Affiliates AECOM SA (Pty) Ltd
Elbroc Mining Products (Pty) Ltd
Namakwa Sands (Pty) Ltd
AEL Mining Services Limited
Engineering and Project Company Ltd
New Concept Mining (Pty) Limited
Air Liquide (PTY) Ltd
eThekwini Municipality
Northam Platinum Ltd - Zondereinde
AMEC Mining and Metals
Exxaro Coal (Pty) Ltd
Osborn Engineered Products SA (Pty) Ltd
AMIRA International Africa (Pty) Ltd
Exxaro Resources Limited
ANDRITZ Delkor(Pty) Ltd
Fasken Martineau
Anglo Operations Ltd
FLSmidth Minerals (Pty) Ltd
Anglo Platinum Management Services (Pty) Ltd
Fluor Daniel SA (Pty) Ltd
Anglogold Ashanti Ltd Atlas Copco Holdings South Africa (Pty) Limited
Outotec (RSA) (Proprietary) Limited PANalytical (Pty) Ltd Paterson and Cooke Consulting Engineers (Pty) Ltd
Franki Africa (Pty) Ltd Johannesburg
Polysius A Division Of Thyssenkrupp Industrial Solutions (Pty) Ltd
Fraser Alexander Group
Precious Metals Refiners Rand Refinery Limited
Glencore
Aurecon South Africa (Pty) Ltd
Redpath Mining (South Africa) (Pty) Ltd
Goba (Pty) Ltd
Rosond (Pty) Ltd
Aveng Moolmans (Pty) Ltd
Hall Core Drilling (Pty) Ltd
Axis House (Pty) Ltd
Hatch (Pty) Ltd
Bafokeng Rasimone Platinum Mine
Herrenknecht AG
Barloworld Equipment -Mining
HPE Hydro Power Equipment (Pty) Ltd
Rustenburg Platinum Mines Limited
BASF Holdings SA (Pty) Ltd
Impala Platinum Limited
SAIEG
Bateman Minerals and Metals (Pty) Ltd
IMS Engineering (Pty) Ltd
Salene Mining (Pty) Ltd
BCL Limited
JENNMAR South Africa
Becker Mining (Pty) Ltd
Joy Global Inc. (Africa)
Sandvik Mining and Construction Delmas (Pty) Ltd
BedRock Mining Support (Pty) Ltd
Leco Africa (Pty) Limited
Sandvik Mining and Construction RSA(Pty) Ltd
Bell Equipment Company (Pty) Ltd
Longyear South Africa (Pty) Ltd
SANIRE
Blue Cube Systems (Pty) Ltd
Lonmin Plc
Sasol Mining(Pty) Ltd
Bluhm Burton Engineering (Pty) Ltd
Ludowici Africa
Scanmin Africa (Pty) Ltd
Blyvooruitzicht Gold Mining Company Ltd
Lull Storm Trading (PTY)Ltd T/A Wekaba Engineering
Sebilo Resources (Pty) Ltd
Magnetech (Pty) Ltd
Senmin International (Pty) Ltd
Magotteaux(PTY) LTD
Shaft Sinkers (Pty) Limited
MBE Minerals SA Pty Ltd
Sibanye Gold (Pty) Ltd
MCC Contracts (Pty) Ltd
Smec SA
MDM Technical Africa (Pty) Ltd
SMS Siemag South Africa (Pty) Ltd
Metalock Industrial Services Africa (Pty)Ltd
SNC Lavalin (Pty) Ltd
Metorex Limited
Sound Mining Solutions (Pty) Ltd
BSC Resources CAE Mining (Pty) Limited Caledonia Mining Corporation CDM Group CGG Services SA Chamber of Mines Concor Mining Concor Technicrete Council for Geoscience Library CSIR-Natural Resources and the Environment
Royal Bafokeng Platinum Roymec Tecvhnologies (Pty) Ltd Runge Pincock Minarco Limited
Metso Minerals (South Africa) (Pty) Ltd Minerals Operations Executive (Pty) Ltd MineRP Holding (Pty) Ltd
SENET
South 32 SRK Consulting SA (Pty) Ltd Technology Innovation Agency Time Mining and Processing (Pty) Ltd
Department of Water Affairs and Forestry
Mintek
Deutsche Securities (Pty) Ltd
MIP Process Technologies
Digby Wells and Associates
Modular Mining Systems Africa (Pty) Ltd
Umgeni Water
Downer EDI Mining
MSA Group (Pty) Ltd
VBKOM Consulting Engineers
DRA Mineral Projects (Pty) Ltd
Multotec (Pty) Ltd
Webber Wentzel
DTP Mining
Murray and Roberts Cementation
Weir Minerals Africa
Duraset
Nalco Africa (Pty) Ltd
WorleyParsons (Pty) Ltd
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Tomra Sorting Solutions Mining (Pty) Ltd Ukwazi Mining Solutions (Pty) Ltd
IP
PONSORSH
EXHIBITS/S
sor ishing to spon Companies w these of y an at t bi and/or exhi contact the events should rdinator -o co ce en confer ssible as soon as po
SAIMM DIARY 2015 or the past 120 years, the Southern African Institute of Mining and Metallurgy, has promoted technical excellence in the minerals industry. We strive to continuously stay at the cutting edge of new developments in the mining and metallurgy industry. The SAIMM acts as the corporate voice for the mining and metallurgy industry in the South African economy. We actively encourage contact and networking between members and the strengthening of ties. The SAIMM offers a variety of conferences that are designed to bring you technical knowledge and information of interest for the good of the industry. Here is a glimpse of the events we have lined up for 2015. Visit our website for more information.
F
For further information contact: Conferencing, SAIMM Pâ&#x20AC;&#x2C6;Oâ&#x20AC;&#x2C6;Box 61127, Marshalltown 2107 Tel: (011) 834-1273/7 Fax: (011) 833-8156 or (011) 838-5923 E-mail: raymond@saimm.co.za
N SYMPOSIUM International Symposium on slope stability in open pit mining and civil engineering 12â&#x20AC;&#x201C;14â&#x20AC;&#x201C; October 2015 In association with the Surface Blasting School 15â&#x20AC;&#x201C;16 October 2015, Cape Town Convention Centre, Cape Town N COLLOQUIUM 13th Annual Southern African Student Colloquim 2015 20 October 2015, Mintek, Randburg, Johannesburg N CONFERENCE Young Professionals 2015 Conference 21â&#x20AC;&#x201C;22 October 2015, Mintek, Randburg, Johannesburg N CONFERENCE AMI: Nuclear Materials Development Network Conference 28â&#x20AC;&#x201C;30 October 2015, Nelson Mandela Metropolitan University, North Campus Conference Centre, Port Elizabeth N SYMPOSIUM MPES 2015: Twenty Third International Symposium on Mine Planning & Equipment Selection 8â&#x20AC;&#x201C;12 November 2015, Sandton Convention Centre, Johannesburg, South Africa
2016 N CONFERENCE Diamonds still Sparkle 2016 Conference 14â&#x20AC;&#x201C;17 March 2016, Gaborone International Convention Centre N CONFERENCE Mine to Market Conference 2016 13â&#x20AC;&#x201C;14 April 2016, South Africa N CONFERENCE The SAMREC/SAMVAL Companion Volume Conference 17â&#x20AC;&#x201C;18 May 2016, Johannesburg N SEMINAR PASTE 2016 International Seminar on Paste and Thickened Tailings May 2016, Kwa-Zulu Natal, South Africa N CONFERENCE 1st International Conference on Solids Handling and Processing A Mineral Processing Perspective 9â&#x20AC;&#x201C;10 June 2016, South Africa N CONFERENCE Hydrometallurgy Conference 2016 1â&#x20AC;&#x201C;3 August 2016, Cape Town N CONFERENCE The Tenth International Heavy Minerals Conference â&#x20AC;&#x2DC;Expanding the horizonâ&#x20AC;&#x2122; 16â&#x20AC;&#x201C;19 August 2016, Sun City, South Africa
Website: http://www.saimm.co.za