（様式 2）
法人番号プロジェクト番号

231029S1201019

平成２４年度～平成２８年度「私立大学戦略的研究基盤形成支援事業」
研究成果報告書概要
１ 学校法人名

名城大学

２ 大学名

名城大学

３ 研究組織名

名城大学・窒化物半導体基盤技術研究センター

４ プロジェクト所在地

名古屋市天白区塩釜口１－５０１ 名城大学天白キャンパス内

５ 研究プロジェクト名

窒化物半導体・新領域エレクトロニクス

６ 研究観点研究拠点を形成する研究（5 年）

７ 研究代表者
研究代表者名

所属部局名

職名

赤﨑 勇

理工学研究科教授

プロジェクト参加研究者数
９ 該当審査区分

７ 名

理工・情報

生物・医歯

人文・社会

１０ 研究プロジェクトに参加する主な研究者
研究者名
赤﨑 勇

上山 智

竹内 哲也

岩谷 素顕

所属・職名

プロジェクトでの研究課題
プロジェクトでの役割
窒化物半導体の結晶成長技
名城大学・術の革新とそれによる革新 窒化物半導体ナノ構造/
理工学研
的な窒化物半導体デバイス 超ヘテロ構造の導入
究科・教授
の実現
ナノコラム結晶成長、ナ
名城大学・
半導体ナノ構造技術と光デ ノ 構 造 の 物 性 評 価 、 高
理工学研
バイスへの応用
性能ナノ発光材料、デバ
究科・教授
イスの実現
名城大学・ 窒化物半導体の結晶成長技 Sb・B を窒化物半導体
理 工 学 研 術の革新とそれによる革新 に添加する技術を開発
究科・准教 的な窒化物半導体デバイス し、革新的なデバイスを
授
の実現
実現する
窒化物半導体のデバイスプ
名城大学・ ロセス技術の革新・貼り合わ
貼り合わせ・基盤剥離技
理 工 学 研 せ・基板剥離技術の確立・ド術やプラズモンを活用す
究科・准教 ーピング技術の確立・電子線る技術を開発する
授
励起レーザ・太陽電池・ディテクターの実現

（様式 2）
法人番号プロジェクト番号（共同研究機関等）
天野 浩

Christian Wetzel

Haiyan Ou

231029S1201019

窒化物半導体の結晶成長技
窒化物半導体ナノ構造/
名古屋大
術の革新とそれによる革新
学・工学研
超ヘテロ構造による的な窒化物半導体デバイス
究科・教授
GaInN・AlGaN の開拓の実現
窒化物半導体の結晶成長技
レンセラー
術の革新とそれによる革新 高 InN モル分率 GaInN
工科大学・的な窒化物半導体デバイス によるデバイスの開拓
教授
の実現
デンマーク
窒化物半導体のデバイスプ
工科大学・
金属プラズモンの応用
ロセス技術の革新
准教授＜研究者の変更状況（研究代表者を含む）＞
旧
ﾌﾟﾛｼﾞｪｸﾄでの研究課題 所属・職名
研究者氏名

プロジェクトでの役割（変更の時期：平成 26 年 4 月 1 日）

新
変更前の所属・職名
ソニー株式会社変更（就任）後の所属・職名 研究者氏名
名城大学・理工学研究科・
宮嶋 孝夫
教授

プロジェクトでの役割
窒化物半導体の物性の解明

（様式 2）
法人番号プロジェクト番号

231029S1201019

１１ 研究の概要（※ 項目全体を１０枚以内で作成）（１）研究プロジェクトの目的・意義及び計画の概要
本研究拠点を形成する目的は、これまで窒化物半導体研究において世界的に高い成果を残してきた名城大学 窒化物半導体基盤技術研究センターを中核に、名古屋大学、デンマー
ク工科大学、およびレンセラー工科大学の研究グループなど世界中の研究機関の技術を結集することによって、新しい結晶成長技術・デバイスプロセス技術を確立する。具体的には、窒化物半導体トンネル接合、窒化物半導体ナノコラム結晶の作製技術、さらにはウェハ接合技術など窒化物半導体において新機能デバイスを実現する上で不可欠な基盤技術の確立をすることを目標に研究を進めた。さらにこれらの基盤技術を基に、最終的に革新的なデバイスの実現し、世界に情報を発信する研究拠点として整備することが本研究課題の目的・意義である。
（２）研究組織
学内：本研究プロジェクトにこれまで参画した研究者は以下の 8 名（学内 5 名、学外 3 名）である
当該研究課題の成研究ﾌﾟﾛｼﾞｪｸﾄにおけ
所属・職
研究者名
果が研究ﾌﾟﾛｼﾞｪｸﾄに
る研究課題
果たす役割
窒化物半導体・新領理工学研究科・教授
赤﨑 勇
研究総括
域エレクトロニクスの実現
窒化物半導体ナノ構
窒化物半導体・新領造の作製・表面プラ
理工学研究科・教授
上山 智
域エレクトロニクスのズモン・高濃度 Si ド
実現ーピング
アンチモン・ボロン添
窒化物半導体・新領理工学研究科・准教
加窒化物半導体の竹内 哲也
域エレクトロニクスの授実現・AlInN 結晶成
実現
長・トンネル接合
貼り合わせ・基板剥離技術の確立・ドー
窒化物半導体・新領理工学研究科・准教
ピング技術の確立・
岩谷 素顕
域エレクトロニクスの授
電子線励起レーザ・
実現
太陽電池・ディテクタ
ーの実現
窒化物半導体の物
窒化物半導体光物理の解明（H26 年度
理工学研究科・教授
宮嶋 孝夫
性の解析と評価
より研究組織に追加）
その他 （特記事項）
H24 から H26 年 3 月まで博士研究員として飯田大輔を雇用し、研究に参加させてきた。主に結晶成長全般およびプラズモンを用いた LED、基板剥離技術等を担当した。H26 年 4 月から
東京理科大学・理学部応用物理学科・助教として栄転・着任し、H29.3 からサウジアラビア・（様式 2）
法人番号プロジェクト番号

231029S1201019

KAIST の Assistant Professor として着任予定である。
さらに、H27 から H29 年 3 月まで、博士研究員として小出典克を雇用し、研究に参画させてきた。
大学院生に関しては本研究グループで以下の通りである。
H24 博士課程学生 3 名 修士課程学生 39 名 （修了 19 名）
H25 博士課程学生 3 名（修了 1 名） 修士課程学生 39 名 （修了 18 名）
H26 博士課程学生 2 名 修士課程学生 35 名 （修了 20 名）
H27 博士課程学生 5 名 修士課程学生 38 名 （修了 17 名）
H28 博士課程学生 5 名（修了 2 名） 修士課程学生 36 名（修了 18 名）
H26 から 2 名の博士課程学生（松井健城・安田俊輝）、さらに H27 から 1 名を追加（小塚祐吾）
を RA として雇用し、本戦略的基盤形成支援事業の取組に従事させている。上記の大学院生
および博士研究員は招待講演 6 件、受賞 12 件などの成果をあげた。
また、研究支援のため秘書を 3 名雇用している（大学予算・科研費・奨学寄附金で執行）
共同研究機関等の連携状況
所属・職

研究者名

名古屋大学・教授

天野 浩

デ ン マ ー ク 工 科 大 Haiyan Ou
学・准教授

レ ン セ ラ ー 工 科 大 Christian Wetzel
学・教授

当該研究課題の成研究ﾌﾟﾛｼﾞｪｸﾄにおけ果が研究ﾌﾟﾛｼﾞｪｸﾄに
る研究課題
果たす役割
窒 化 物 ナ ノ ワ イ ヤ ・ 窒化物半導体・新領
量子ドット・原子層エ 域エレクトロニクスの
ピ タ キ シ ー の 成 長 ・ 実現
評価・デバイス応用、電子線励起レーザや太陽電池、分極を用いた紫外 LED の開発等で連携表面プラズモンの理 窒化物半導体・新領論 計 算 と 作 製 、 PD 域エレクトロニクスのの飯田大輔を派遣、 実現
博士課程学生を 2 名（各 1~2 か月）受け入れ
緑色領域の LED 作 窒化物半導体・新領製、太陽電池の作製 域エレクトロニクスの実現（３）研究施設・設備等（研究施設）◆クリーンルーム 150m2＋155m2、デバイス・光学評価室 55m2 使用者: 4 名（研究装置）◆有機金属化合物気相成長装置(①、④、⑩、29 に相当) 4 台（利用時間: 1 日
12 時間）◆フォトリソグラフィ装置(⑧、⑪)＋電子ビーム露光装置(30)（利用時間: 1 日 12 時
間）◆ICP エッチング 装置(⑤)（利用時間: 1 日 12 時間）◆エキシマレーザ(31)（利用時間:
1 日 12 時間）◆EB 蒸着装置(28)（利用時間: 1 日 12 時間） ◆スパッタリング装置(32)（利用時間: 1 日 12 時間）◆カソードルミネッセンス装置付き走査電子顕微鏡(2,15)（利用時間:

（様式 2）
法人番号プロジェクト番号

231029S1201019

1 日 5 時間）◆FIB・マイクロサンプリング機能付き透過電子顕微鏡（利用時間: 1 日 3 時間）(33)◆X 線回析装置(34)（利用時間: 1 日 12 時間）◆フォトルミネッセンス装置（利用時間:
1 日 3 時間）◆レーザ顕微鏡装置(35)（利用時間: 1 日 3 時間）◆半導体デバイス評価装置(14、27、26)（利用時間: 1 日 3 時間）◆結晶評価装置(3、12、17、21、25)（利用時間: 1 日
12 時間）◆デバイスシミュレータ（SiLENSe、TCAD 等）（利用時間: 1 日 12 時間）
および本予算で導入したウェハー接合装置(9)（H24）（利用時間: 1 日 12 時間）、半導体原料供給設備(22)（利用時間: 1 日 12 時間）、昇華法装置（H25）(7)（利用時間: 1 日 6 時間）（４）研究成果の概要 ※下記、１３及び１４に対応する成果には下線及び＊を付すこと。
＜現在までの進捗状況及び達成度＞
本研究グループは、窒化物半導体材料を中核とし、これまでに世界で初めて同材料の高品
質単結晶の実現、伝導性制御などを達成し、世界に先駆けて窒化物半導体による青色/紫色
LED をはじめ、世界最短波長半導体レーザ、高効率紫外 LED、超高速トランジスタなど、フロンティアデバイスを実現してきた。本研究プロジェクトでは、これまで構築してきた研究基盤を基に結晶成長技術およびデバイスプロセス技術を革新し、最終的に革新的なデバイスの実現を目指すことを目指している。結晶成長・デバイスプロセスにおいて以下のような成果を残した。
1. 有機金属化合物気相成長法におけるその場観察 X 線回折測定法の確立を行い、それをGaN および GaInN/GaN ヘテロ接合に適用する技術を確立(*1 に対応)
2. アンチモンを窒化物半導体に 0.4%取り込むことに成功(*2 に対応)
3. AlInN の高速・高品質結晶成長技術の確立(*3 に対応)
4. デバイス適用可能な超低抵抗 n 型窒化物半導体(5.9×10-4Ωcm)の実現(*4 に対応)
5. 低抵抗窒化物半導体トンネル接合およびそれをデバイスに適用(*5 に対応)
6. バルク AlN 基板上へのホモエピタキシャル成長技術の確立(*6 に対応)
7. 紫外/可視 LED における分極の理解(*7 に対応)
8. ITO 電極を用いた基板貼り合わせ技術を確立(*8 に対応)
9. AlGaInP と窒化物半導体の直接接合技術を確立(*9 に対応)
10. GaN 基板からのレーザ剥離技術を確立(*10 に対応)
11. ナノインプリントを用いたプラズモン適用技術を確立(*11 に対応)
12. AlGaInP と窒化物半導体を ITO を用いたウェハ接合技術の確立
13. ナノコラム結晶・量子殻構造の作製技術の確立
14. 周期利得構造を目指した電子・正孔注入効率を制御する技術を確立
さらに、これらの結晶成長技術/デバイスプロセスを基に以下のような革新的なデバイスを実現した。
① 窒化物半導体・超高感度・可視光/紫外線ディテクターの実現（受光感度 10４A/W 以上）の実現(*12 に対応)
② 窒化物半導体において世界最高効率の変換効率を持つ太陽電池の実現(*13 に対応)
③ GaN/AlGaN 系においては世界初となる電子線励起のレーザの実現(*14 に対応)
④ 反射率 99.9%以上の AlInN 半導体多層膜反射鏡の実現(*15 に対応)
⑤ キャリア注入制御 2 波長で発光する窒化物半導体可視 LED の実現(*16 に対応)
⑥ 電流ドループが極めて少ない近紫外 LED の実現(*17 に対応)
⑦ トンネル接合を用いた電流狭窄・マイクロ LED の実現(*18 に対応)
⑧ 窒化物半導体トンネル接合を用いた多接合窒化物半導体太陽電池の実現(*19 に対応)
⑨ 光出力 3mW（世界最高出力：H29.3 時点）の窒化物半導体面発光レーザを実現⑩ 窒化物半導体面発光レーザに於いて、DBR 層に電流を流す縦伝導デバイスを実現⑪ 周期利得構造を用いたデバイス構造において、窒化物半導体面発光レーザを実現（様式 2）
法人番号プロジェクト番号

231029S1201019

上記のように、この約 5 年間で予定していた研究内容に加え、新たなるデバイスを実現するなどの成果を残してきた。
＜優れた成果が上がった点＞
上記に記載した研究成果のうち特に優れた研究成果としては、①有機金属化合物気相成長法におけるその場観察 X 線回折測定法の確立を行い、それを GaN、GaInN/GaN ヘテロ接合、AlGaN/GaN 接合、GaInN/GaN 超格子構造に適用する技術を確立②窒化物半導体・超高感度・可視光/紫外線ディテクターの実現、③GaN においては世界初となる電子線励起のレーザの実現④キャリア注入を制御することによって 2 波長で発光する LED の実現⑤低抵抗
窒化物半導体トンネル接合の実現⑥窒化物半導体太陽電池において世界最高効率⑥分極
ドーピングを紫外 LED に適用⑦世界最高出力の窒化物半導体系面発光レーザを実現があげられる。以下に本成果の概説をする。
① 有機金属化合物気相成長法におけるその場観察 X 線回折測定法の確立を行い、それをGaN および GaInN/GaN ヘテロ接合に適用する技術を確立
窒化物半導体デバイスのほとんどは有機金属化合物気相成長法で作製されている。その結晶成長技術を極限まで追求する上で、その場観察技術は極めて重要である。これまでに、光を用いた、膜厚・基板の反り・温度等をその場観察する技術が広く報告されており、実際にデバイス製造に用いられている。本研究ではさらに一歩進め、X 線を用いたその場観察技術を検討した。これまでの検討で、光では観測できなかった GaInN の緩和過程を正確にモニターすることが可能であることを報告するなど極めて優れた成果を残した。これらの成果は 5 件の招待講演、学生の受賞 1 件など対外的に高い評価を得ている。
② 窒化物半導体・超高感度・可視光/紫外線ディテクターの実現
窒化物半導体のバンドギャップエネルギーは紫外から可視全域、赤外領域にわたっており、発光素子だけでなく受光素子としての応用も期待されている。これまで、本グループをはじめとした世界中の研究機関から pin フォトダイオードや光導電セルなど様々なデバイスが報告されてきたが、受光感度が低いことなど実用化には程遠い物しか得られていなかった。本研究
課題では、その問題点を解決するために、AlGaN/GaN ヘテロ接合を用いた 2 次元電子ガスとpn 接合によって形成された空乏層を活用する構造を新たに開発し、それによって極めて高い
受光感度が得られることを明らかにした。本デバイスの性能は、既存の窒化物半導体を用いた受光素子の性能を凌駕するだけではなく、Si を用いた pin フォトダイオードやアバランシェフ
ォトダイオードの受光感度を 3～4 桁程度凌駕しており、受光素子としては最も受光感度が高い光電子増倍管と同程度の物が実現可能であることを明らかにした。また、受光波長はAlGaN の組成や p 型層に GaInN を用いることによって制御可能であることも明らかにした。これらのデバイスは、紫外 LED の実用化にともないキーデバイスとして発展していくことが期待される。これらの成果は招待講演が 4 件、学生の受賞が 2 件に選出されるなど高い評価を得ている。
③窒化物半導体系においては世界初となる電子線励起のレーザの実現
窒化物半導体レーザは、これまでに青紫色・青色・緑色レーザが実用化されているが、物性的には紫外領域や可視全域、さらには赤外領域に亙っている。このうち紫外領域は他の半導体材料では実現困難な領域であり、医療や工業用途向けに実現が強く期待されている。
本研究グループでは、世界最短波長紫外レーザの実現（2007 年：当時）をするなど世界に先駆けた成果を多数残してきたが、さらなる短波長化は現状の技術の延長線では困難である。
その最も大きな理由は、高い正孔濃度を持つ p 型 AlGaN を実現することが困難であり、そのためレーザ発振に必要な電流注入ができないということである。その問題点を解決するためには他のキャリア注入の手法を確立することが重要であり、本グループでは電子線励起に注目した。電子線励起は電流注入の問題がなくなり、紫外レーザの実現が可能になると期待さ（様式 2）
法人番号プロジェクト番号

231029S1201019

れるが、これまで窒化物半導体材料において電子線励起によるレーザの報告はなかった。
本グループでは、名古屋大学の天野教授さらにはスタンレー電気との共同研究を実施し、その問題点の解決を検討し、世界ではじめて電子線励起による窒化物半導体レーザの実現を確認した。今後は、同手法を紫外領域のデバイスに応用し、最終的には波長が制御可能な
紫外線レーザの実現を目指す予定である。これらの結果は招待講演が 1 件など学術的に高い評価を得ている。
④キャリア注入を制御することによって 2 波長で発光する LED の実現
一般に LED では一つの活性層を一つの pn 接合で挟み込むため、一つの波長でしか発光しない。本素子では、二つ以上の活性層を挟みこんで多波長で発光する LED の実現を目論んだ。この場合、各活性層への電子と正孔（キャリア）の均一な注入が必要である。本研究グループでは、キャリアの注入状態に着目し、その制御により 2 波長で発光することが可能である
ことを見出した。特に中間層と呼ぶ層を形成し、そこに適量の Mg を添加することによって、効率よく 2 つの波長で発光させることが可能になった。本成果は、学生が受賞（1 件）するなど高く評価され、今後低コスト白色光源や青色面発光レーザなどへの適用も期待できる。
⑤低抵抗窒化物半導体トンネル接合の実現
窒化物半導体トンネル接合は多接合太陽電池や面発光レーザの基盤技術として重要であるが、これまでの報告ではトンネル接合部の抵抗が非常に高い（～50Ω）という問題点が存在した。本グループでは p 型 GaInN の適用、活性化処理の手法の確立などを行い、抵抗が 5Ω程度という一桁低いトンネル接合を実現した。また、この技術を電流狭窄構造や多接合太陽電池に適用し、世界で初めてのトンネル接合を用いた多接合太陽電池を実現するなどの成果を残した。これらの結果は招待講演が 3 件など学術的に高い評価を得ている。
⑥ 窒化物半導体太陽電池において世界最高効率を実現
窒化物半導体太陽電池において、超格子構造・透明電極・集光器等を用いることによって窒化物半導体太陽電池において世界最高効率を実現した。これらの成果は招待講演 5 件など学会で高い評価を得ている。
＜課題となった点＞
上記のように、本研究課題では窒化物半導体の結晶成長をはじめとした基盤技術の確立を狙い、トンネル接合や分極ドーピング、貼りあわせ技術などこれまで窒化物半導体材料ではデバイス応用が進められていない多くの技術を確立した。その一方で、社会実装を考えると、これを 1 歩進めて革新的な光デバイスへの適用が不可欠であると考えられる。
＜自己評価の実施結果と対応状況＞
本プロジェクトでは、多数の外部発表や年 1 回のシンポジウムを主催し、外部の有識者からの評価・アドバイスをうけながら研究を研究を推進した。
＜外部（第三者）評価の実施結果と対応状況＞
中間評価において学部の学識研究者による中間評価を実施し、それに基づいて適切な対応を行った。具体的には、研究拠点の創出に向けた取り組みを強化し、窒化物半導体研究
に関する拠点形成を目指した。
＜研究期間終了後の展望＞
上述のように、本研究課題では、窒化物半導体の結晶成長やウェハーボンディング技術など新規の基盤技術を形成した。これらの基盤技術は新しい光デバイス創製に極めて有効であると考えられ、今後本グループでは新たなる光デバイスの創製に向け検討を進めていく予定である。

（様式 2）
法人番号プロジェクト番号

231029S1201019＜研究成果の副次的効果＞
本研究課題の研究代表者の赤﨑勇および外部の共同研究者である天野浩は高効率青色
LED の発明による成果で 2014 年のノーベル物理学賞を受賞した。また、それをきっかけに、
名古屋大学を中心に GaN 研究コンソーシアムを形成し、本拠点も参画している。さらに、本研究を基に派生した技術を活用し、2016 年度より名城大学光デバイス研究センターを開設し、本センターを中心に次世代の光デバイスの開発に取り掛かっている。さらに、本研究室の学生が以下のような受賞、招待講演に選ばれるなど研究拠点を中心に社会的に高い評価を得ている。
［学生の受賞］
1. 窒化物半導体研究会 発表奨励賞 (2012) ：石黒真未
2. 応用物理学会 講演奨励賞 (2012）： 石黒真未
3. 窒化物半導体研究会 研究奨励賞（2013）： 飯田大輔
4. 窒化物半導体研究会 発表奨励賞（2014）：池山和希
5. 窒化物半導体研究会 発表奨励賞（2014）：小塚祐吾
6. 窒化物半導体研究会 発表奨励賞（2014）：松井健城
7. 窒化物半導体研究会 発表奨励賞（2014）：安田俊輝
8. 窒化物半導体研究会 発表奨励賞（2013）： 山本泰司
9. ISGN Best Young Scientist Award (2015): A. Yoshikawa
10. ISGN Best Young Scientist Award (2015): K. Takeda
11. ISGN Best Young Scientist Award (2015): D. Komori
12. ISGN Best Young Scientist Award (2015): T. Furuta
13. ISGN Young Scientist Award (2015): K. Matsui
14. 応用物理学会 講演奨励賞（2015）: 池山和希
15. LEDIA Best Student Award (2016): K. Matsui
16. 窒化物半導体結晶成長講演会 発表奨励賞 (2016): 吉川陽
17. 窒化物半導体結晶成長講演会 発表奨励賞 (2016): 牛田彩希
18. 応用物理学会 講演奨励賞 (2016): 吉川陽［学生の基調講演・招待講演リスト］
1. IWN (2012 札幌): M. Ishiguro
2. 応用物理学会（2012 東京）: 石黒真未
3. ACCGE(2013 アメリカ): T. Yamamoto
4. ICMOVPE (2014 スイス）: T. Yamamoto
5. ISGN (2015 浜松): K. Takeda
6. ECCG (2015 イタリア): J. Ohsumi
7. 応用物理学会（2016 新潟） 吉川陽
8. EMN meeting（2016 ハンガリー) K. Matsui
9. EMN meeting (2016 ハンガリー) A. Yoshikawa
10. 応用物理学会（2016 東京） 池山和希
11. ISPlasma（2017、名古屋） 林 菜摘
１２ キーワード（当該研究内容をよく表していると思われるものを８項目以内で記載してください。）（１） 窒化物半導体（２） 結晶成長（３） その場観察技術（４） トンネル接合（５） LED（６） レーザ（様式 2）
法人番号プロジェクト番号（７） ディテクター

231029S1201019（８） 太陽電池

１３ 研究発表の状況（研究論文等公表状況。印刷中も含む。）
上記、１１(４)に記載した研究成果に対応するものには＊を付すこと。
＜雑誌論文＞
論文名、著者名、掲載誌名、査読の有無、巻、最初と最後の頁、発表年（西暦）について記入
してください（左記の各項目が網羅されていれば、項目の順序を入れ替えても可）。また、現在から発表年次順に遡り、通し番号を付してください。
1. "Laser lift-off of AlN/sapphire for UV light-emitting diodes" Hiroki Aoshima, Kenichiro
Takeda, Kosuke Takehara, Shun Ito, Mikiko Mori, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: physica status solidi (c)Vol. 9 (2012) 753.
2. "Growth of GaN and AlGaN on (100) β-Ga2O3 substrates" Shun Ito, Kenichiro
Takeda, Kengo Nagata, Hiroki Aoshima, Kosuke Takehara, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki and Hiroshi Amano: physica statussolidi (c) Vol. 9 (2012) 519.
3. "Small current collapse in AlGaN/GaN HFETs on a-plane GaN self-standing
substrate" T. Sugiyama, Y. Honda, M. Yamaguchi, H. Amano, Y. Isobe, M. Iwaya, T.
Takeuchi, S. Kamiyama, I. Akasaki, M. Imade, Y. Kitaoka and Y. Mori: physica
status solidi (c) Vol. 9 (2012) 875.
4. "Fabrication of AlInN/AlN/GaInN/GaN heterostructure field-effect transistors"
Kazuya Ikeda, Yasuhiro Isobe, Hiromichi Ikki, Tatsuyuki Sakakibara, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki and Hiroshi Amano: physica
status solidi (c) Vol. 9 (2012) 942.
5. "Enhancement of two-dimensional electron gases in AlGaN-channel
high-electron-mobility transistors with AlN barrier layers" Shin Hashimoto, Katsushi
Akita, Yoshiyuki Yamamoto, Masaki Ueno, Takao Nakamura, Kenichiro Takeda,
Motoaki Iwaya, Yoshio Honda, and Hiroshi Amano: physica status solidi (a) Vol.
209 (2012) 501
6. ＊ "Properties of nitride-based photovoltaic cells under concentrated lightillumination" Shota Yamamoto, Mikiko Mori, Yosuke Kuwahara, Takahiro Fujii,
Tatsuo Nakao, Shinichiro Kondo, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki and Hiroshi Amano: physica status solidi (RRL) - Rapid
Research Letters, Vol. 6 (2012) 145.(*13 に対応)
7. "Epitaxial Indium Tin Oxide Film Deposited on Sapphire Substrate by Solid-Source
Electron Cyclotron Resonance Plasma" Satoru Kaneko, Hironori Torii, Masayasu
Soga, Kensuke Akiyama, Motoaki Iwaya, Mamoru Yoshimoto, and Takao
Amazawa: Japanese Journal of Applied Physics Vol. 51 (2012) 01AC02.
8. "Indium-Tin Oxide/Al Reflective Electrodes for Ultraviolet Light-Emitting Diodes"
Kosuke Takehara, Kenichiro Takeda, Shun Ito, Hiroki Aoshima, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano:
Japanese Journal of Applied Physics Vol. 51 (2012) 042101.
9. "Crack-Free AlN/GaN Distributed Bragg Reflectors on AlN Templates" Kouta Yagi,
Mitsuru Kaga, Kouji Yamashita, Kenichirou Takeda, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama, Hiroshi Amano, and Isamu Akasaki: Japanese
Journal of Applied Physics 51 (2012) 051001.
10. ＊"Correlation between Device Performance and Defects in GaInN-Based Solar
Cells" Mikiko Mori, Shinichiro Kondo, Shota Yamamoto, Tatsuro Nakao, Takahiro
Fujii, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and（様式 2）
法人番号プロジェクト番号

231029S1201019

Hiroshi Amano: Applied Physics Express 5 (2012) 082301.(*13 に対応)
11. "White light-emitting diode based on fluorescent SiC" Satoshi Kamiyama, Motoaki
Iwaya, Tetsuya Takeuchi, Isamu Akasaki, Rositza Yakimova, and Mikael Syväjärvi:Thin Solid Films 522 (2012) 23-25.
12. “MOVPE growth of nonpolar a-plane GaN with low oxygen contamination andspecular surface on a freestanding GaN substrate” Yasuhiro Isobe, Hiromichi Ikki,
Tatsuyuki Sakakibara, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu
Akasaki, Takayuki Sugiyama, Hiroshi Amano, Mamoru Imade, Yusuke Mori: Journalof Crystal Growth Vol. 351 (2012) pp. 126-130.
13. “Development of AlN/diamond heterojunction field effect transistors” Masataka
Imura, Ryoma Hayakawa, Hirotaka Ohsato, Eiichiro Watanabe, Daiju Tsuya,
Takahiro Nagata, Meiyong Liao, Yasuo Koide, Jun-ichi Yamamoto, Kazuhito Ban,
Motoaki Iwaya, Hiroshi Amano: Diamond and Related Materials Vol. 24 (2012) pp.
206-209.
14. ＊“In situ X-ray diffraction monitoring during metalorganic vapor phase epitaxy
growth of low-temperature-GaN buffer layer” Daisuke Iida, Mihoko Sowa, Yasunari
Kondo, Daiki Tanaka, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, andIsamu Akasaki: Journal of Crystal Growth 361 (2012) 1.(*1 に対応)
15. ＊ "Nitride-based hetero-field-effect transistor-type photosensors with extremely
high photosensitivity" Mami Ishiguro, Kazuya Ikeda, Masataka Mizuno, Motoaki
Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: Physica Status
Solidi RRL 7 (2013) 215–217.(*12 に対応)
16. ＊"Analysis of strain relaxation process in GaInN/GaN heterostructure by in situ
X-ray diffraction monitoring during metalorganic vapor-phase epitaxial growth"
Daisuke Iida, Yasunari Kondo, Mihoko Sowa, Toru Sugiyama, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: Physica Status Solidi
RRL 7 (2013) 211–214.(*1 に対応)
17. “Trench-Shaped Defects on AlGaInN Quantum Wells Grown under Different Growth
Pressures” Tomoyuki Suzuki, Mitsuru Kaga, Kouichi Naniwae, Tsukasa Kitano,
Keisuke Hirano, Tetsuya Takeuchi, Satoshi Kamiyama, Motoaki Iwaya, and Isamu
Akasaki: Japanese Journal of Applied Physics 52 (2013) 08JB27.
18. “ControloftheDetection
Wavelengthin
AlGaN/GaN-Based
Hetero-Field-Effect-Transistor Photosensors” Mami Ishiguro, Kazuya Ikeda,
Masataka Mizuno, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, andIsamu Akasaki: Japanese Journal of Applied Physics 52 (2013) 08JF02.
19. ＊“Carrier Injections in Nitride-Based Light Emitting Diodes Including Two Active
Regions with Mg-Doped Intermediate Layers” Kenjo Matsui, Koji Yamashita,
Mitsuru Kaga, Takatoshi Morita, Tomoyuki Suzuki, Tetsuya Takeuch, Satoshi
Kamiyama, Motoaki Iwaya, and Isamu Akasaki: Japanese Journal of Applied
Physics 52 (2013) 08JG02.(*16 に対応)
20. “Combination of Indium–Tin Oxide and SiO2/AlN Dielectric Multilayer Reflective
Electrodes for Ultraviolet-Light-Emitting Diodes” Tsubasa Nakashima, Kenichiro
Takeda, Hiroshi Shinzato, Motoaki Iwaya, Satoshi Kamiyama, Tetsuya Takeuchi,
Isamu Akasaki, and Hiroshi Amano: Japanese Journal of Applied Physics 52 (2013)
08JG07.
21. ＊ “Concentrating Properties of Nitride-Based Solar Cells Using Different
Electrodes” Mikiko Mori, Shinichiro Kondo, Shota Yamamoto, Tatsuro Nakao,
Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi
Amano: Japanese Journal of Applied Physics 52 (2013) 08JH02.(*13 に対応)
22. ＊“GaInN-Based Tunnel Junctions in n–p–n Light Emitting Diodes” Mitsuru Kaga,（様式 2）
法人番号プロジェクト番号

231029S1201019

Takatoshi Morita, Yuka Kuwano, Kouji Yamashita, Kouta Yagi, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: Japanese Journal ofApplied Physics 52 (2013) 08JH07.(*18 に対応)
23. ＊“Investigations of Polarization-Induced Hole Accumulations and Vertical Hole
Conductions in GaN/AlGaN Heterostructures” Toshiki Yasuda, Kouta Yagi,
Tomoyuki Suzuki, Tsubasa Nakashima, Masahiro Watanabe, Tetsuya Takeuchi,
Motoaki Iwaya, Satoshi Kamiyama, and Isamu Akasaki: Japanese Journal ofApplied Physics 52 (2013) 08JJ05.(*7 に対応)
24. “Study on Efficiency Component Estimation of 405 nm Light Emitting Diodes from
Electroluminescence and Photoluminescence Intensities” Kazuki Aoyama, Atsushi
Suzuki, Tsukasa Kitano, Satoshi Kamiyama, Tetsuya Takeuchi, Motoaki Iwaya, andIsamu Akasaki: Japanese Journal of Applied Physics 52 (2013) 08JL16.
25. "Dislocation density dependence of stimulated emission characteristics in AlGaN/Al
multiquantum wells" Yuko Matsubara, Tomoaki Yamada, Kenichiro Takeda, Motoaki
Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki and Hiroshi Amano:
physica status solidi (c) Volume 10 (2013) pp. 1537–1540.
26. ＊ “In situ X-ray diffraction monitoring of GaInN/GaN super lattice duringorganometalic vapor phase epitaxy growth” Journal of Crystal Growth 393 (2014)
108.(*1 に対応)
27. ＊ “Extremely Low-Resistivity and High-Carrier-Concentration Si-Doped
Al0.05Ga0.95N Toru Sugiyama, Daisuke Iida, Toshiki Yasuda, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama and Isamu Akasaki: Applied Physics Express Vol. 6(2013) 121002.(*4 に対応)
28. ＊ "Homoepitaxial growth of AlN layers on freestanding AlN substrate bymetalorganic vapor phase epitaxy" Tomohiro Morishita, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki: Journal of Crystal Growth Volume
390 (2014) pp. 46-50.(*6 に対応)
29. ＊ “Multijunction GaInN-based solar cells using a tunnel junction” Hironori
Kurokawa, Mitsuru Kaga, Tomomi Goda, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki and Hiroshi Amano: Appllied Physics Express Vol. 7(2014) 034104.(*20 に対応)
30. ＊ “Control of growth mode in Mg-doped GaN/AlN heterostructure” Tomohiro
Morishita, Kosuke Sato, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama andIsamu Akasaki: Japanese Journal of Applied Phyics Vol. 53 (2014) 030305.(*6 に対応)
31. ＊“Control of crystallinity of GaN grown on sapphire substrate by metalorganic
vapor phase epitaxy using in situ X-raydiffraction monitoring method” Journal of
Crystal Growth 401 (2014) 367–371.(*1 および*13 に対応)
32. “Laser lift-off technique for freestanding GaN substrate using an In droplet formedby thermal decomposition of GaInN and its application to light-emitting diodes”
Daisuke Iida, Syunsuke Kawai, Nobuaki Ema, Takayoshi Tsuchiya, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: Applied Physics Letters
105 (2014) 072101.
33. 竹内哲也、岩谷素顕、上山 智、赤﨑 勇、「窒化物半導体面発光レーザーの現状」、
照明学会誌 第 100 巻第 5 号（2016）

（様式 2）
法人番号プロジェクト番号

231029S1201019

34. Matsui, K., Kozuka, Y., Ikeyama, K., Horikawa, K., Furuta, T., Akagi, T., ... &Akasaki, I. (2016). GaN-based vertical cavity surface emitting lasers with periodic
gain structures. Japanese Journal of Applied Physics, 55(5S), 05FJ08.
35. Takeda, K., Iwaya, M., Takeuchi, T., Kamiyama, S., & Akasaki, I. (2016). Electricalproperties of n-type AlGaN with high Si concentration. Japanese Journal of Applied
Physics, 55(5S), 05FE02.
36. Komori, D., Takarabe, K., Takeuchi, T., Miyajima, T., Kamiyama, S., Iwaya, M., &Akasaki, I. (2016). GaNSb alloys grown with H2 and N2 carrier gases. Japanese
Journal of Applied Physics, 55(5S), 05FD01.
37. Osumi, J., Ishihara, K., Yamamoto, T., Iwaya, M., Takeuchi, T., Kamiyama, S., &Akasaki, I. (2016). GaN barrier layer dependence of critical thickness in GaInN/GaN
superlattice on GaN characterized by in situ X-ray diffraction. Japanese Journal ofApplied Physics, 55(5S), 05FD11.
38. Furuta, T., Matsui, K., Horikawa, K., Ikeyama, K., Kozuka, Y., Yoshida, S., ... &Akasaki, I. (2016). Room-temperature CW operation of a nitride-based
vertical-cavity surface-emitting laser using thick GaInN quantum wells. Japanese
Journal of Applied Physics, 55(5S), 05FJ11.
39. Yamamoto, Y., Yoshikawa, A., Kusafuka, T., Okumura, T., Iwaya, M., Takeuchi, T., ...
&Akasaki,I.
(2016).
Realizationofhigh-performance
hetero-field-effect-transistor-type ultraviolet photosensors using p-type GaN
comprising three-dimensional island crystals. Japanese Journal of Applied Physics,
55(5S), 05FJ07.
40. Yoshikawa, A., Yamamoto, Y., Murase, T., Iwaya, M., Takeuchi, T., Kamiyama, S., &Akasaki,I.
(2016).
High-photosensitivity
AlGaN-based
UV
heterostructure-field-effect-transistor-type photosensors. Japanese Journal ofApplied Physics, 55(5S), 05FJ04.
41. Mori, K., Takeda, K., Kusafuka, T., Iwaya, M., Takeuchi, T., Kamiyama, S., ... &Amano, H. (2016). Low-ohmic-contact-resistance V-based electrode for n-type
AlGaN with high AlN molar fraction. Japanese Journal of Applied Physics, 55(5S),
05FL03.
42. Yoshida, S., Ikeyama, K., Yasuda, T., Furuta, T., Takeuchi, T., Iwaya, M., ... &Akasaki, I. (2016). Electron and hole accumulations at GaN/AlInN/GaN interfaces
and conductive n-type AlInN/GaN distributed Bragg reflectors. Japanese Journal ofApplied Physics, 55(5S), 05FD10.
43. Takasuka, D., Akatsuka, Y., Ino, M., Koide, N., Takeuchi, T., Iwaya, M., ... & Akasaki,
I. (2016). GaInN-based tunnel junctions with graded layers. Applied Physics
Express, 9(8), 081005.
44. Ikeyama, K., Kozuka, Y., Matsui, K., Yoshida, S., Akagi, T., Akatsuka, Y., ... &Akasaki, I. (2016). Room-temperature continuous-wave operation of GaN-based
vertical-cavity surface-emitting lasers with n-type conducting AlInN/GaN distributed
Bragg reflectors. Applied Physics Express, 9(10), 102101.
＜図書＞

（様式 2）
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1. “青い光に魅せられて 青色 LED 開発物語” 赤崎 勇、日本経済新聞出版社、233 ペー
ジ、（2013）
2. “ワイドギャップ半導体―あけぼのから最前線へ” 赤﨑勇、松波 弘之、培風館 404 ページ（2013）. （6.2 紫外線発光ダイオード応用（12 ページ）を岩谷素顕が担当
3. ”ポストシリコン半導体 －ナノ成膜ダイナミクスと基板・界面効果－”（2.1 窒化物半導体の特徴とデバイス動向）、竹内哲也、エヌ・ティー・エス、11 ページ（2013）.
＜学会発表＞
15. ＊"窒化物半導体太陽電池の現状とその将来展望" 岩谷素顕、竹内哲也、上山智、赤﨑勇、天野浩：日本結晶成長学会ナノ構造・エピタキシャル成長分科会シンポジウム「高効率ナノエピタキシャル太陽電池の最先端」、東京大学、2012.4.26-27. [特別講演](*13 に対応)
16. "膜厚変化に伴う GaInN/GaN ヘテロ接合の微細構造観察" 松原大幸、飯田大輔、 杉
山 徹、近藤保成、曽和美保子、岩谷素顕、竹内哲也、上山智、赤﨑勇: 第 4 回 窒化物
半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28.
17. ＊"n 型 GaN 直下におけるトンネル接合および p 型 GaN 活性化の検討" 桑野侑香、山下浩司、加賀充、森田隆敏、竹内哲也、岩谷素顕、上山智、赤﨑勇: 第 4 回 窒化物半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28. (*5 に対応)
18. "埋め込み型 GaN ナノコラム結晶の MOVPE 成長に関する検討" 梅田慎也、加藤嵩裕、
北野司、近藤俊行、松原大幸、上山智、竹内哲也、岩谷素顕、赤﨑勇: 第 4 回 窒化物
半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28.
19. "GaInN 量子井戸への Al 添加による AlGaInN 量子井戸の作製と評価" 鈴木智行、 加
賀充、北野司、難波江宏一、平野敬祐、竹内哲也、上山智、岩谷素、赤﨑勇: 第 4 回
窒化物半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28.
20. ＊"X 線その場観察 MOVPE により評価した GaInN/GaN の臨界膜厚" 近藤保成、飯田
大輔、杉山徹、曽和美保子、松原大幸、岩谷素顕、竹内哲也、上山智、赤﨑勇: 第 4 回
窒化物半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28. (*1 に対応)
21. ＊"中間層を用いた窒化物半導体二波長発光ダイオードにおけるキャリア注入の検討"
松井健城、山下浩司、加賀充、森田隆敏、鈴木智行、竹内哲也、上山智、岩谷素顕、赤﨑 勇 : 第 4 回 窒 化 物 半 導 体 結 晶 成 長 講 演 会 ( プ レ IWN2012) 、 東 京 大 学 、
2012.4.27-28. (*16 に対応)
22. ＊"Ⅲ族窒化物半導体を用いた低抵抗トンネル接合" 森田隆敏、加賀充、桑野侑香、山下浩司、松井健城、竹内哲也、上山智、岩谷素顕、赤﨑勇: 第 4 回 窒化物半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28. (*5 に対応)
23. ＊"p 型 GaInN を用いた高感度な FET 型光センサーに関する研究" 石黒真未、池田和弥、水野正隆、岩谷素顕、竹内哲也、上山智、赤﨑勇: 第 4 回 窒化物半導体結晶成長
講演会 (プレ IWN2012)、東京大学、2012.4.27-28. (*12 に対応)
24. "非極性面上Ⅲ族窒化物半導体電界効果トランジスタに関する研究" 水野正隆、池田和弥、石黒真未、松原大幸、岩谷素顕、竹内哲也、上山智、天野浩、赤﨑勇: 第 4 回 窒化物半導体結晶成長講演会 (プレ IWN2012)、東京大学、2012.4.27-28.
25. ＊"Optimizations of Nitride Semiconductor-Based Tunnel Junctions" Mitsuru Kaga,
Kouji Yamashita, Takatoshi Morita, Yuka Kuwano, Kouta Yagi, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: 16th International
Conference on Metal Organic Vapor Phase Epitaxy (ICMOVPE-XVI), Pusan, Korea,
2012. 5.20-25. (*5 に対応)
26. ＊ "In-Situ Real-Time X-Ray Diffraction Study during Ga1-xInxN/GaN

（様式 2）
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231029S1201019

Heterostructure Growth by MOVPE for Understanding Strain Relaxation
Mechanism" Motoaki Iwaya, Daisuke Iida, Daiki Tanaka, Toru Sugiyama, Mihoko
Sowa, Yasunari Kondo, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki:
16th International Conference on Metal Organic Vapor Phase Epitaxy(ICMOVPE-XVI), Pusan, Korea, 2012. 5.20-25. (*1 に対応)
27. "Microstructure of AlGaN on Low-Dislocation Density AlN Underlying Layer Grown
by Epitaxial Lateral Overgrowth" Kimiyasu Ide, Junichi Yamamoto, Motoaki Iwaya,
Satoshi Kamiyama, Tetsuya Takeuchi, Isamu Akasaki, and Hiroshi Amano: 16th
International Conference on Metal Organic Vapor Phase Epitaxy (ICMOVPE-XVI),
Pusan, Korea, 2012. 5.20-25.
28. "Investigation of AlN/GaN Multilayer Stacks for DBR Applications" Kouta Yagi,
Mitsuru Kaga, Kouji Yamashita, Kenichirou Takeda, Motoaki Iwaya, Tetsuya
Takeichi, Satoshi Kamiyama, Hiroshi Amano, and Isamu Akasaki: 16th International
Conference on Metal Organic Vapor Phase Epitaxy (ICMOVPE-XVI), Pusan, Korea,
2012. 5.20-25.
29. ＊"基板・半導体層のナノ加工と半導体デバイスへの応用" 岩谷素顕: 日本機械学会
情報・知能・精密機器部門 分科会 第１回 「窒化物半導体デバイスに関わる超精密加工プロセス研究会」発会記念 特別講演会、九州大学、2012 年 5 月 30 日. [招待講演](*1 に対応)30. "窒化物半導体エピタキシャル成長の新展開" 天野浩、本田善央、山口雅史、今出完、
森勇介、岩谷素顕：応用物理学会 第 137 回結晶工学分科会研究会「窒化物半導体光デバイスの最前線-基板・エピ成長と評価技術-」京都テレサ、2012 年 6 月 15 日. [特別講演]31. ＊"Observation of GaInN strain relaxation by in situ X-ray diffraction monitoring
during metalorganic vapour phase epitaxy growth" M. Iwaya, D. Iida, T. Sugiyama,
M. Sowa, Y. Kondo, H. Matsubara, T. Takeuchi, S. Kamiyama, and I. Akasaki: 4thInternational Symposium on Growth of III-Nitrides, St. Petersburg, Russia,
2012.7.16-7.19. (*1 に対応)
32. ＊"Composition dependence of critical thicknesses in GaInN/GaN characterized byin situ X-ray diffraction measurement", Y. Kondo, D. Iida, T. Sugiyama, M. Sowa, H.
Matsubara, M. Iwaya, T. Takeuchi, S.Kamiyama, and I. Akasaki: 4th International
Symposium on Growth of III-Nitrides, St. Petersburg, Russia, 2012.7.16-7.19. (*1 に対応)
33. ＊"Strain Relaxation Mechanism in GaInN/GaN Heterostructure Characterized by insitu X-ray Diffraction Monitoring During Metalorganic Vapor Phase Epitaxy Growth",
Motoaki Iwaya, Yasunari Kondo, Hiroyuki Matsubara, Mihoko Sowa, Toru
Sugiyama, Daisuke Iida, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki:
39th International Symposium on Compound Semiconductors, University ofCalifornia Santa Barbara, CA USA, August 27-30, 2012. (*1 に対応)
34. "非極性 AlGaN/GaN HFET の AlGaN バリアの AlN モル分率および Si 濃度依存性" 水野正隆，磯部康裕，池田和弥，石黒真未，岩谷素顕，竹内哲也，上山 智，赤崎 勇，天野 浩: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日.
35. ＊"ピエゾ・自発分極電荷への正孔蓄積の検討" 安田俊輝，矢木康太，鈴木智行，中嶋翼，渡邉雅大，竹内哲也，上山 智，岩谷素顕，赤崎 勇: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日. (*7 に対応)
36. ＊"GaInN/GaN ヘテロ接合におけるミスフィット転位の観察" 松原大幸，飯田大輔，杉山徹，近藤保成，曽和美保子，梅田慎也，岩谷素顕，竹内哲也，上山 智，赤崎 勇: 第73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日. (*1 に対応)

（様式 2）
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37. "AlGaN/AlN における下地 AlN 転位密度依存性" 井手公康，松原由布子，岩谷素顕，
竹内哲也，上山 智，赤崎 勇，天野 浩: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日.
38. "ITO と SiO2/AlN 誘電体多層膜を組み合わせた電極による 350nm 紫外 LED の高効率化"中嶋 翼，竹田健一郎，新里紘史，岩谷素顕，上山 智，竹内哲也，赤崎 勇，天野浩: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日.
39. ＊"窒化物半導体を用いた高感度な HFET 型光センサー"石黒真未，池田和弥，水野正隆，岩谷素顕，竹内哲也，上山 智，赤崎 勇: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日. (*12 に対応)
40. ＊"窒化物半導体太陽電池の集光特性における電極構造の最適化" 森美貴子，山本翔太，近藤真一郎，中尾達郎，岩谷素顕，竹内哲也，上山 智，赤崎 勇，天野 浩: 第 73
回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日. (*13 に対応)
41. ＊"Mg ドープ中間層を用いた二波長発光ダイオードにおける電流注入依存性量の検討"
松井健城，山下浩司，加賀 充，森田隆敏，鈴木智行，竹内哲也，上山 智，岩谷素顕，
赤崎 勇: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14
日. (*16 に対応)
42. ＊"III 族窒化物半導体トンネル接合を用いた n-p 接合 LED" 森田隆敏，加賀 充，桑野侑香，松井健城，竹内哲也，上山 智，岩谷素顕，赤崎 勇: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日. (*5 に対応)
43. "100 チャンネル 10μm 径マイクロ LED アレイの作製" 渡邉雅大，山下浩司，加賀 充，
鈴木智行，森田隆敏，竹内哲也，上山 智，岩谷素顕，赤崎 勇: 第 73 回応用物理学会講演会、愛媛大学・松山大学、2012 年 9 月 11 日-14 日.
44. "EL と PL を組み合わせた LED 効率成分の導出方法の検討" 青山和樹，鈴木敦志，北野 司，上山 智，竹内哲也，岩谷素顕，赤崎 勇: 第 73 回応用物理学会講演会、愛媛
大学・松山大学、2012 年 9 月 11 日-14 日.
45. ＊"Realization of high efficiency nitride-based solar cells" Motoaki Iwaya, Mikiko
Mori, Shinichiro Kondo, Shota Yamamoto, Tatsuro Nakao, Takahiro Fujii, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: The International
Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
[Invited Talk] (*13 に対応)
46. ＊"Optimization of crystalline quality of GaN using low temperature buffer layer byin situ X-ray diffraction monitoring" Daiki Tanaka, Daisuke Iida, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: The International
Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
(*1 に対応)
47. ＊"AlN mole fraction dependence of polarization induced hole concentrations in
GaN/AlGaN heterostructures" Toshiki Yasuda, Kouta Yagi, Tomoyuki Suzuki,
Tsubasa Nakashima, Masahiro Watanabe, Tetsuya Takeuchi, Satoshi Kamiyama,Motoaki Iwaya, and Isamu Akasaki: The International Workshop on Nitride
Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19. (*7 に対応)
48. "Fabrication of AlInN/GaInN MIS heterostructure field-effect transistors" Kazuya
Ikeda, Yasuhiro Isobe, Hiromichi Ikki, Kimiyasu Ide, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: The International
Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
49. ＊ "High sensitivity extremely nitride-based heterofield-effect-transistor-type
photosensors" Mami Ishiguro, Kazuya Ikeda, Masataka Mizuno, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: The International（様式 2）
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231029S1201019

Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
[Invited Talk] (*13 に対応)
50. ＊"Composition dependence of critical thicknesses in GaInN/GaN characterized byin situ X-ray diffraction measurement" Yasunari Kondo, Daisuke Iida, Toru
Sugiyama, Hiroyuki Matsubara, Mihoko Sowa, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: The International Workshop on Nitride
Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19. (*1 に対応)
51. ＊"Strain relaxation mechanism in GaInN/GaN heterostructure characterized by insitu XRD monitoring during growth and ex situ measurements" Hiroyuki Matsubara,
Daisuke Iida, Toru Sugiyama, Yasunari Kondo, Mihoko Sowa, Shinya Umeda,
Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: The
International Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan,
2012.10.14-19. (*1 に対応)
52. "Fabrication and characterization of AlGaInN quantum wells" Tomoyuki Suzuki,
Mitsuru Kaga, Koichi Naniwae, Tsukasa Kitano, Kesuke Hirano, Tetsuya Takeuchi,
Satoshi Kamiyama, Motoaki Iwaya, and Isamu Akasaki: The International
Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
53. "Study on efficiency component estimation by electroluminescence andphotoluminescence intensities" Kazuki Aoyama, Atsushi Suzuki, Tsukasa Kitano,
Satoshi Kamiyama, Tetsuya Takeuchi, Motoaki Iwaya, and Isamu Akasaki: The
International Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan,
2012.10.14-19.
54. "Nonpolar a-plane AlGaN/GaN heterostructure field-effect transistor grown on
freestanding GaN substrate" Masataka Mizuno, Yasuhiro Isobe, Kazuya Ikeda,
Mami Isiguro, Motoaki Iwaya, Tetuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki,
and Hiroshi Amano: The International Workshop on Nitride Semiconductors(IWN2012), Sapporo, Japan, 2012.10.14-19.
55. "Combination of ITO and SiO2/AlN dielectric multilayer reflective electrodes for UV
light-emitting diodes" Tsubasa Nakashima, Kenichirou Takeda, Yuko Matsubara,Motoaki Iwaya, Tetyuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi
Amano: The International Workshop on Nitride Semiconductors (IWN2012),
Sapporo, Japan, 2012.10.14-19.
56. ＊"Carrier injections in nitride-based light emitting diodes including two active
regions with Mg-doped intermediate layers" Kenjo Matsui, Koji Yamashita, Mitsuru
Kaga, Takatoshi Morita, Tomoyuki Suzuki, Tetsuya Takeuchi, Satoshi Kamiyama,Motoaki Iwaya, and Isamu Akasaki: The International Workshop on Nitride
Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19. (*1 に対応)
57. ＊"Concentrating properties of nitride-based solar cells using different electrodes"
Mikiko Mori, Shota Yamamoto, Shinichiro Kondo, Tatsuro Nakao, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: The
International Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan,
2012.10.14-19. (*13 に対応)
58. ＊"GaInN-based tunnel junctions in n-p-n light emitting diodes" Mitsuru Kaga,
Takatoshi Morita, Yuka Kuwano, Kouji Yamashita, Kouta Yagi, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: The International
Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
(*5 に対応)59. "Microstructure analysis of AlGaN on AlN underlying layers with different threading
dislocation densities" Kimiyasu Ide, Yuko Matsubara, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: The International（様式 2）
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Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
60. "MOVPE growth of embedded GaN nanocolumn" Shinya Umeda, Takahiro Kato,
Tsukasa Kitano, Toshiyuki Kondo, Hiroyuki Matsubara, Satoshi Kamiyama, Tetsuya
Takeuchi, Motoaki Iwaya, and Isamu Akasaki: The International Workshop on
Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
61. ＊"Lateral hydrogen diffusion at p-GaN layers in nitride-based LEDs with tunnel
junctions" Yuka Kuwano, Mitsuru Kaga, Takatoshi Morita, Tetsuya Takeuchi,
Motoaki Iwaya, Satoshi Kamiyama, and Isamu Akasaki: The International
Workshop on Nitride Semiconductors (IWN2012), Sapporo, Japan, 2012.10.14-19.
(*5 に対応)62. "Carrier profiling technique for DUV LEDs" Yuta Furusawa, Inazu Tetsuhiko, Shinya
Fukahori, Cyril Pernot, Myonghee Kim, Takehiko Fujita, Yosuke Nagasawa, Akira
Hirano, Masamichi Ippommatsu, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Yoshio Honda, Masashi Yamaguchi, Hiroshi Amano, and Isamu
Akasaki: The International Workshop on Nitride Semiconductors (IWN2012),
Sapporo, Japan, 2012.10.14-19.
63. ＊"窒化物半導体太陽電池の高効率化" 岩谷素顕、竹内哲也、上山智、赤﨑勇、天野浩: 第 42 回結晶成長国内会議、九州大学、2012 年 11 月 9 日-11 日. [招待講演] (*13に対応)
64. "ドループ改善に向けた窒化物半導体 np 構造 LED の作製" 森田隆敏、加賀 充、桑野侑香、松井健城、竹内哲也、上山 智、岩谷素顕、赤﨑勇：電子情報通信学会 ED、
CPM、LQE 合同研究会、大阪市立大学、2012 年 11 月 29 日-30 日.
65. "サファイア基板上へのモスアイ構造の形成方法の検討および LED への応用"土屋貴義、梅田慎也、曽和美保子、近藤俊行、北野 司、森 みどり、鈴木敦志、難波江宏一、
関根 均、岩谷素顕、竹内哲也、上山智、赤﨑勇：電子情報通信学会 ED、CPM、LQE
合同研究会、大阪市立大学、2012 年 11 月 29 日-30 日.
66. ＊"n-GaN 表面層を有する構造内 p-GaN の Mg アクセプタ活性化" 桑野侑香、加賀充、森田隆敏、山下浩司、南川大智、竹内哲也、岩谷素顕、上山 智、赤﨑勇：電子情報通信学会 ED、CPM、LQE 合同研究会、大阪市立大学、2012 年 11 月 29 日-30 日.
(*5 に対応)
67. ＊“Observation of group III nitride semiconductors by in situ X-ray diffraction
monitoring during metalorganic vapor phase epitaxy growth” Motoaki Iwaya,
Daisuke Iida, Mihoko Sowa, Yasunari Kondo, Daiki Tanaka, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: 2012 Collaborative Conference on Crystal
Growth, Orland, FL, USA, 2012.12.11-14. [Invited Talk] (*1 に対応)
68. ＊"High-sensitivity HFET type photosensors with a p-GaInN gate" Mami Ishiguro,
Kazuya Ikeda, Masataka Mizuno, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki: SPIE Photonic west 2013, San Francisco, USA,
2013.2.2-8. (*12 に対応)
69. "Influence of growth interruption on performance of nitride-based blue LED" Kazuki
Aoyama, Atsushi Suzuki, Tsukasa Kitano, Naoki Sone, Satoshi Kamiyama, Tetsuya
Takeuchi, Motoaki Iwaya, Isamu Akasaki: SPIE Photonic west 2013, San Francisco,USA, 2013.2.2-8.
70. "Introduction of the Moth-eye patterned sapphire substrate technology for costeffective high-performance LED" Koichi Naniwae, Midori Mori, Toshiyuki Kondo,
Atsushi Suzuki, Tsukasa Kitano, Satoshi Kamiyama, Motoaki Iwaya, Tetsuya
Takeuchi, Isamu Akasaki: SPIE Photonic west 2013, San Francisco, USA,
2013.2.2-8. [Invited Talk]
71. "Fabrication of moth-eye patterned sapphire substrate (MPSS) and influence of

（様式 2）
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height of corns on the performance of blue LEDs on MPSS" Takayoshi Tsuchiya,
Shinya Umeda, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu
Akasaki, Toshiyuki Kondo, Tsukasa Kitano, Midori Mori, Atsushi Suzuki, Fumiharu
Teramae, Hitoshi Sekine: SPIE Photonic west 2013, San Francisco, USA,
2013.2.2-8.
72. "Performance of nitride-based light emitting diodes using an Indium-zinc-oxidetransparent electrode with moth-eye structure" Shugo Mizutani, Satoshi
Nakashima, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki,
Toshiyuki Kondo, Fumiharu Teramae, Atsushi Suzuki, Tsukasa Kitano, Midori Mori,
Masahito Matsubara: SPIE Photonic west 2013, San Francisco, USA, 2013.2.2-8.
73. ＊"GaInN トンネル接合素子の低抵抗化" 森田隆敏、加賀 充、桑野侑香、松井健城、
渡邉雅大、竹内哲也、岩谷素顕、上山 智、赤﨑 勇: 第 60 回応用物理学会春季学術
講演会、神奈川工科大学、2013 年 3 月 26 日-30 日. (*5 に対応)
74. ＊"GaN 低温成長における Sb 添加の効果" 鈴木智行、笹島浩希、松原由布子、加賀充、竹内哲也、上山 智、岩谷素顕、赤﨑 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日. (*2 に対応)
75. ＊"LED 構造上 GaN トンネル接合の低抵抗化" 南川大智、加賀 充、岩谷素顕、竹内哲也、上山 智、赤﨑 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、
2013 年 3 月 26 日-30 日. (*7 に対応)
76. ＊"非対称 AlN/GaN 多層膜反射鏡の設計と作製" 萩原康大、矢木康太、安田俊輝、
竹田健一郎、竹内哲也、岩谷素顕、上山 智、赤﨑 勇: 第 60 回応用物理学会春季学
術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日. (*15 に対応)
77. "ナノ ELO による m 面 GaInN 厚膜の検討" 小﨑桂矢、近藤真一郎、土屋貴義、岩谷素顕、上山 智、竹内哲也、赤﨑 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日.
78. ＊"MOVPE 法による AlInN 高速成長の影響" 小塚祐吾、鈴木智行、森田隆敏、竹内
哲也、上山 智、岩谷素顕、赤﨑 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日. (*3 に対応)
79. ＊"MOVPE により作製した GaInN/GaN 超格子構造のその場観察 X 線回折法による評価" 山本泰司、飯田大輔、近藤保成、曽和美保子、岩谷素顕、竹内哲也、上山 智、
赤崎 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26
日-30 日. (*1 に対応)
80. ＊"表面プラズモンカップリングを利用した GaInN/GaN 量子井戸の評価" 飯田大輔、
Yuntian Chen、Yiyu Ou、Ahmed Fadil、Oleksii Kopylov、岩谷素顕、竹内哲也、上山智、赤﨑 勇、Haiyan Ou: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、
2013 年 3 月 26 日-30 日. (*11 に対応)
81. ＊"窒化物半導体 HFET 型光センサー構造の最適化" 石黒真未、池田和弥、水野正隆、岩谷素顕、上山 智、竹内哲也、赤崎 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日.［受賞記念講演］ (*12 に対応)
82. ＊"GaInN 超格子構造最適化による窒化物半導体太陽電池の高性能化" 黒川泰視、
近藤真一郎、山本翔太、森美貴子、岩谷素顕、竹内哲也、上山 智。赤﨑 勇、天野浩: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30
日. (*13 に対応)
83. "350nm 紫外 LED の光取り出し効率の向上" 中嶋 翼、竹田健一郎、新里紘史、岩谷素顕、上山 智、竹内哲也、赤﨑 勇、天野 浩: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日.
84. "モスアイ加工サファイア基板による LED の光取り出し効率の改善"土屋貴義、梅田慎（様式 2）
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也、曽和美保子、河合俊介、近藤俊行、北野 司、森みどり、鈴木敦志、赤﨑 勇、関根均、難波江宏一、岩谷素顕、竹内哲也、上山 智: 第 60 回応用物理学会春季学術講
演会、神奈川工科大学、2013 年 3 月 26 日-30 日.
85. "電子ブロック層上活性層による電子オーバーフローの直接観測": 林 健人、松井健
城、森田隆敏、鈴木智行、竹内哲也、上山 智、岩谷素顕、赤﨑 勇: 第 60 回応用物
理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日.
86. "横方向 p 型活性化における雰囲気ガス依存性" 桑野侑香、加賀 充、森田隆敏、山下浩司、南川大智、竹内哲也、岩谷素顕、上山 智、赤﨑 勇: 第 60 回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日.
87. "窒化物半導体二波長発光ダイオードにおける発光強度比の安定性"松井健城、山下浩
司、加賀 充、森田隆敏、鈴木智行、竹内哲也、上山 智、岩谷素顕、赤﨑 勇: 第 60
回応用物理学会春季学術講演会、神奈川工科大学、2013 年 3 月 26 日-30 日.
88. "Direct Evidence of Electron Overflow by Monitoring Emissions from Second Active
Region in Nitride-Based Blue LEDs" K. Hayashi, K. Matsui , T. Morita, T. Suzuki, T.
Takeuchi, S. Kamiyama, M. Iwaya, and I. Akasaki: Conference on LED and Its
Industrial Application ’13, Pacifico Yokohama, Yokohama, Japan, 2013.4.23-25.
89. ＊"Nitride-Based p-Side Down LEDs on Tunnel Junction" T. Morita, M. Kaga, Y.
Kuwano, K. Matsui, M. Watanabe, T. Takeuchi, S. Kamiyama, M. Iwaya, and I.
Akasaki: Conference on LED and Its Industrial Application ’13, Pacifico Yokohama,
Yokohama, Japan, 2013.4.23-25. (*5 に対応)
90. ＊"Observation of GaInN/GaN Superlattice Structures by In Situ X-ray Diffraction
Monitoring during Metalorganic Vapor-Phase Epitaxial Growth" T. Yamamoto, D.
Iida, Y. Kondo, M. Sowa, S. Umeda, T. Kato, M. Iwaya, T. Takeuchi, S. Kamiyama,
and I. Akasaki: Conference on LED and Its Industrial Application ’13, Pacifico
Yokohama, Yokohama, Japan, 2013.4.23-25. (*1 および*13 に対応)
91. "Threshold Power Density Reduction in AlGaN/AlN Multiquntum Wells DUV (288nm) Optical Pumped Laser" T. Yamada, Y. Matsubara, H. Shinzato, K. Takeda, M.
Iwaya, T. Takeuchi, S. Kamiyama, I. Akasaki, and H. Amano: Conference on LED
and Its Industrial Application ’13, Pacifico Yokohama, Yokohama, Japan,
2013.4.23-25.
92. ＊“Factors for obtaining a high open-circuit voltage in GaInN-based solar cells”
Motoaki Iwaya, Shinichiro Kondo, Shota Yamamoto, Mikiko Mori, Yasushi
Kurokawa, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, Hiroshi Amano:
Asia-Pacific Workshop on Widegap Semiconductors 2013, Tamsui, New Taipei City,
Taiwan, May 12-15, 2013. (*1 に対応)
93. ＊ “Bandgap dependence in nitride semiconductor-based tunnel junctions”, D.
Minamikawa, M. Kaga, Y. Kuwano, T.Morita, T. Takeuchi, S. Kamiyama, M. Iwaya, I.
Akasaki: Asia-Pacific Workshop on Widegap Semiconductors 2013, Tamsui, New
Taipei City, Taiwan, May 12-15, 2013. (*5 に対応)
94. ＊“High sensitivity group III nitride semiconductor based heterostructure field effecttransistor type photosensors” Motoaki Iwaya, Mami Ishiguro, Masataka Mizuno,
Kazuya Ikeda, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki: The
Collaborative Conference on Crystal Growth (3CG), Cuncun, Mexico, June 10-13,
2013. [invited Talk] (*12 に対応)
95. ＊“窒化物半導体における分極の影響と発光素子への応用” 竹内哲也、岩谷素顕、上山智、赤﨑勇: 第 5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日
-22 日. [チュートリアル講演] (*7 に対応)
96. “端面コーティングによる光励起 DUV レーザの閾値励起強度密度低減について” 山田（様式 2）
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知明, 松原由布子, 新里紘史, 竹田健一郎, 岩谷素顕, 竹内哲也, 上山智, 赤崎勇,
天野浩: 第 5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日.
97. “AlGaN/AlN 多重量子井戸の誘導放出における転位密度依存性” 松原 由布子,山田知明, 中嶋 翼, 竹田 健一郎, 岩谷 素顕, 竹内 哲也, 上山 智, 赤﨑 勇, 天野 浩:
第 5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日.
98. ＊“GaInN における InN 組成揺らぎの改善による窒化物太陽電池の高効率化” 黒川泰視, 近藤真一郎, 山本翔太, 森美貴子, 岩谷素顕, 竹内哲也, 上山智, 赤﨑 勇,天野浩: 第 5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日. (*13に対応)
99. ＊“その場観察 X 線回折法による GaInN/GaN 超格子構造の評価” 山本泰司, 飯田大輔, 近藤保成, 曽和美保子, 梅田慎也, 岩谷素顕, 竹内哲也, 上山智, 赤崎勇: 第5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日. (*13 に対応)
100.
＊“表面プラズモンカップリングを利用した緑色 LED の内部量子効率の向上” 飯田
大輔、Yuntian Chen、Yiyu Ou、Ahmed Fadil、Oleksii Kopylov、岩谷素顕、竹内哲也、
上山智、赤﨑勇、Haiyan Ou: 第 5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日.( (*11 に対応)
101.
＊“トンネル接合を有する青色マイクロ LED の電流電圧特性” 中島啓介, 渡邉雅
大 , 加賀 充 , 鈴木智行 , 南川大智, 竹内哲也, 上山 智 , 岩谷素顕, 赤崎 勇: 第5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日. (*5 に対応)
102.
＊“新規 GaNSb 混晶の作製と GaSb モル分率成長温度依存性” 笹島浩希, 鈴木智行, 松原由布子, 竹内哲也, 上山智, 岩谷素顕, 赤崎勇: 第 5 回 窒化物半導体結晶成長講演会、大阪大学、2013 年 6 月 21 日-22 日. (*2 に対応)
103.
“ ＊ In situ X-ray diffraction monitoring of OMVPE GaInN/GaN superlattice
growth” Taiji Yamamoto, Daisuke Iida, Yasunari Kondo, Mihoko Sowa, Shinya
Umeda, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki: The
19th American Conference on Crystal Growth and Epitaxy, Keystone, Colorado,
USA, July 21 - 26, 2013. (*1 に対応)
104.
＊“In situ X-ray diffraction monitoring of GaInN growth by metalorganic vapor
phase epitaxy” Motoaki Iwaya, Taiji Yamamoto, Daisuke Iida, Yasunari Kondo,
Mihoko Sowa, Shinya Umeda, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu
Akasaki: 17th International Conference on Crystal Growth and Epitaxy, Warsaw,
Poland, August 11-16, 2013. (*1 に対応)
105.
＊ “Extreme Low-Resistivity and High-Carrier-Concentration Si-Doped
Al0.05Ga0.95N” Toru Sugiyama, Daisuke Iida, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: 10th International Conference of Nitride
Semiconductors, Washington DC, USA, August 25-30, 2013. (*4 に対応)
106.
＊“Realization of the High Conversion Efficiency Solar Cells using Nitride
Semiconductors” Motoaki Iwaya, Yasushi Kurokawa, Yosuke Katsu, Taiji
Yamamoto, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi
Amano: 10th International Conference of Nitride Semiconductors, Washington DC,
USA, August 25-30, 2013. [Invited Talk] (*13,*8, *9, *19 に対応)
107.
＊ “GaNSb Alloys Grown by Low Temperature Metalorganie Vapor Phase
Epitaxy” Tornoyuki Suzuki, Hiroki Sasajima, Yuuko Matsubara, Yugo Kozuka,
Tetsuya Takeuchi, Satoshi Karniyama, Motoaki Iwaya, and Isamu Akasaki: 10th
International Conference of Nitride Semiconductors, Washington DC, USA, August25-30, 2013. (*1 に対応)
108.
＊“Obtaining a High Open-Circuit Voltage in GaInN-Based Solar Cells” Hironori

（様式 2）
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Kurokawa, Shinichiro Kondo, Mihoko Mori, Motoaki Iwaya, Tetsuya Takeuchi,
Isamu Akasaki, and Hiroshi Amano: 10th International Conference of Nitride
Semiconductors, Washington DC, USA, August 25-30, 2013. (*13 に対応)
109.
＊“Stable Balance of Emission Intensities from Two Active Regions in Nitride
Semiconductor-Based Light Emitting Diodes” Kenjo Matsui, Koji Yamashita, Mitsuru
Kaga, Takatoshi Morita, Yuka Kuwano, Tomoyuki Suzuki, Tetsuya Takeuchi,
Satoshi Kamiyama, Motoaki Iwaya, and Isamu Akasaki: 10th International
Conference of Nitride Semiconductors, Washington DC, USA, August 25-30, 2013.
(*16 に対応)
110.
“＊10-μm-Square Micro LED Array with Tunnel Junction” Masahiro Watanabe,
Mitsuru Kaga, Koji Yamashita, Tomoyuki Suzuki, Daichi Minamikawa, Yuka
Kuwano, Tetsuya Takeuchi, Satoshi Karniyarna, Motoaki Iwaya, and Isarnu
Akasaki: 10th International Conference of Nitride Semiconductors, Washington DC,
USA, August 25-30, 2013. (*18 に対応)
111.
“Advantages of the Moth-Eye Patterned Sapphire Substrate for the High
Perfortnance Nitride Based LEDs” Toshiyuki Kondo, Tsukasa Kitano, Atsushi
Suzuki, Midori Mori, Koichi Naniwae, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, and Isamu Akasaki: 10th International Conference of Nitride
Semiconductors, Washington DC, USA, August 25-30, 2013.
112.
“Improvement of Light Extraction Efficiency in 350-nm Emission UVLight-Emitting Diodes” Tsubasa Nakashima, Kenitirou Takeda, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: 10th
International Conference of Nitride Semiconductors, Washington DC, USA, August25-30, 2013.
113.
＊"トンネル接合を活用した窒化物半導体多接合太陽電池の作製"黒川泰視，合田智美，加賀充，岩谷素顕，竹内哲也，上山智，赤﨑勇，天野浩: 第 74 回応用物理学会
秋季学術講演会, 同志社大学, 2013 年 9 月 15 日-20 日. (*13 および*19 に対応)
114.
＊"GaNSb 混晶 GaSb モル分率の Sb/N 比依存性" 笹島 浩希，鈴木 智行，松原
由布子，竹内 哲也，上山 智，岩谷 素顕，赤﨑 勇: 第 74 回応用物理学会秋季学術講演会, 同志社大学, 2013 年 9 月 15 日-20 日. (*2 に対応)
115.
＊"MOVPE 法による AlInN の高速成長" 小塚 祐吾，安田 俊輝，鈴木 智行，竹内哲也，上山 智，岩谷 素顕，赤﨑 勇: 第 74 回応用物理学会秋季学術講演会, 同志社
大学, 2013 年 9 月 15 日-20 日. (*3 に対応)
116.
＊"青色面発光レーザに向けた電流狭窄構造の検討" 中島 啓介，渡邉 雅大，加賀 充，鈴木 智行，南川 大智，竹内 哲也，上山 智，岩谷 素顕，赤﨑 勇: 第 74 回応用物理学会秋季学術講演会, 同志社大学, 2013 年 9 月 15 日-20 日. (*18 に対応)
117.
＊"窒化物半導体を用いた HFET 型光センサーの暗電流検討" 石黒 真未，池田和弥，水野 正隆，村瀬 卓弥，岩谷 素顕，上山 智，赤崎 勇: 第 74 回応用物理学会
秋季学術講演会, 同志社大学, 2013 年 9 月 15 日-20 日. (*12 に対応)
118.
" VSL 法を用いた AlGaN/AlN 多重量子井戸の光学利得測定" 山田 知明，松原 由
布子，竹田 健一郎，岩谷 素顕，竹内 哲也，上山 智，赤崎 勇，天野 浩: 第 74 回応用物理学会秋季学術講演会, 同志社大学, 2013 年 9 月 15 日-20 日.
119.
"AlGaN/AlN 多重量子井戸の光学利得の転位密度依存性" 松原 由布子，山田 知
明，竹田 健一郎，岩谷 素顕，竹内 哲也，上山 智，赤﨑 勇，天野 浩: 第 74 回応用物理学会秋季学術講演会, 同志社大学, 2013 年 9 月 15 日-20 日.
120.
＊“In situ X-ray diffraction monitoring of GaInN/GaN superlattices growth bymetalorganic vapor phase epitaxy” Taiji Yamamoto, Daisuke Iida, Yasunari Kondo ,Mihoko Sowa , Shinya Umeda , Motoaki Iwaya, Tetsuya Takeuchi, Satoshi（様式 2）
法人番号プロジェクト番号

231029S1201019

Kamiyama, and Isamu Akasaki: 2013 JSAP-MRS Joint Symposia, Doshisha Univ.,
Kyoto, Japan, September 18-19, 2013. (*1 に対応)
121.
＊ “Enhanced internal quantum efficiency of green emission GaInN/GaN
multiple quantum wells by surface plasmon coupling” Daisuke Iida, Yuntian Chen,
Yiyu Ou, Ahmed Fadil, Oleksii Kopylov, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki, and Haiyan Ou: 2013 JSAP-MRS Joint Symposia,
Doshisha Univ., Kyoto, Japan, September 18-19, 2013. (*11 に対応)
122.
“Reduction of the threshold power density in AlGaN/AlN multiquntum wells DUV(288 nm) optical pumped lasers” Tomoaki Yamada, Yuko Matsubara, Hiroshi
Shinzato, Kenichiro Takeda, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama,
Isamu Akasaki, and Hiroshi Amano: 2013 JSAP-MRS Joint Symposia, Doshisha
Univ., Kyoto, Japan, September 18-19, 2013.
123.
“Quantitative evaluation of electron overflow by monitoring emissions fromsecond active region in nitride-based blue LEDs” Kento Hayashi, Kenjo Matsui,
Takatoshi Morita, Tomoyuki Suzuki, Tetsuya Takeuchi, Satoshi Kamiyama, Motoaki
Iwaya, and Isamu Akasaki: 2013 JSAP-MRS Joint Symposia, Doshisha Univ.,
Kyoto, Japan, September 18-19, 2013.
124.
“10 um square micro LED and 100 channel array with tunnel junctions”
Masahiro Watanabe, Mitsuru Kaga, Koji Yamashita, Tomoyuki Suzuki, Daichi
Minamikawa, Yuka Kuwano, Tetsuya Takeuchi, Satoshi Kamiyama, Motoaki Iwaya,
and Isamu Akasaki: 2013 JSAP-MRS Joint Symposia, Doshisha Univ., Kyoto,
Japan, September 18-19, 2013.
125.
“Asymmetric AlN/GaN-DBRs with high reflectivity” Koudai Hagiwara, Toshiki
Yasuda, Kouta Yagi, Kenichirou Takeda, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, and Isamu Akasaki: 2013 JSAP-MRS Joint Symposia, Doshisha Univ.,
Kyoto, Japan, September 18-19, 2013.
126.
“Moth-eye Patterned Sapphire Substrate technology for cost effective highperformance LED” Toshiyuki Kondo, Tsukasa Kitano, Atsushi Suzuki, Midori Mori,
Koichi Naniwae, Satoshi Kamiyama, Tetsuya Takeuchi , Motoaki Iwaya, and Isamu
Akasaki: 2013 JSAP-MRS Joint Symposia, Doshisha Univ., Kyoto, Japan,
September 18-19, 2013.
127.
“MOVPE growth of embedded GaN nanocolumn” Shinya Umeda, Takahiro
Kato, Tsukasa Kitano, Toshiyuki Kondo, Satoshi Kamiyama, Tetsuya Takeuchi,
Motoaki Iwaya, and Isamu Akasaki: 2013 JSAP-MRS Joint Symposia, Doshisha
Univ., Kyoto, Japan, September 18-19, 2013.
128.
＊ "Externally high sensitivity group III nitride semiconductor based
heterostructure field effect transistor type photosensors" M. Ishiguro, M. Mizuno, M.
Iwaya, T. Takeuchi, S. Kamiyama, I. Akasaki: 3rd International Conference on
Materials and Applications for Sensors and Transducers, Prague, Czech Republic,
September 13-17, 2013. (*12 に対応)
129.
＊"オーバーフロー抑制に向けた np 接合 GaInN-LED" 竹内哲也, 岩谷素顕, 上山 智, 赤﨑 勇: 日本学術振興会ワイドギャップ半導体光・電子デバイス第 162 委員
会 第 86 回研究会, 名古屋, 2013 年 10 月 10 日. [依頼講演] (*7, *17 に対応)
130.
"Dislocation density dependence of stimulated emission characteristics in
AlGaN/AlN multi-quantum wells" Motoaki Iwaya, Yuko Matsubara, Tomoaki
Yamada, Kenichiro Takeda, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasak,and Hiroshi Amano: 8th International Workshop on Bulk Nitride Semiconductor,
Kloster Seeon, Germany, September 30-October 5, 2013.
131.
"近接昇華法による SiC 結晶のドーピング濃度均一化に関する検討" 畔柳有孝，前田智彦， 赤澤絵理，江間伸明，寺前文晴，藤元直樹，難波江宏一 2，上山智，竹内哲（様式 2）
法人番号プロジェクト番号

231029S1201019也，岩谷素顕，赤﨑勇: 第 43 回結晶成長国内会議, 長野,2013 年 11 月 6 日-8 日.
132.
＊"新規 GaNSb 混晶の作製及び高 GaSb モル分率に向けた検討"笹島浩希，鈴木智行，松原由布子，竹内哲也，上山智，岩谷素顕，赤崎勇: 第 43 回結晶成長国内会
議, 長野,2013 年 11 月 6 日-8 日. (*2 に対応)
133.
＊"AlInN の高速成長に関する研究" 小塚祐吾，安田俊輝， 鈴木智行，竹内哲也，
上山智，岩谷素顕，赤﨑勇: 第 43 回結晶成長国内会議, 長野,2013 年 11 月 6 日-8 日.
(*3 に対応)134.
"細線加工 m 面 GaN 基板上への GaInN 成長" 小﨑桂矢, 近藤真一郎, 梅田慎也,
岩谷素顕, 竹内哲也, 上山智, 赤﨑勇: 第五回薄膜太陽電池セミナープログラム，名古屋大学，2013 年 11 月 14 日-15 日.
135.
＊"in-situX 線を用いた GaInN/GaN 薄膜超格子構造の観察" 山本泰司, 飯田大輔, 近藤保成, 曽和美保子, 梅田慎也 , 松原大幸, 岩谷素顕, 竹内哲也, 上山智,
赤﨑勇: 第五回薄膜太陽電池セミナープログラム，名古屋大学，2013 年 11 月 14 日-15
日. (*1 に対応)
136.
＊"トンネル接合を用いた III 族窒化物半導体多接合型太陽電池"合田智美, 黒川泰視，森田隆敏, 岩谷素顕, 竹内哲也, 上山智, 赤﨑勇, 天野浩: 応用物理学会 SC 東海地区学術講演会 2013，名古屋大学，2013 年 11 月 16 日. (*19 に対応)
137.
"モスアイ加工サファイア基板を用いた窒化物系 LED の性能向上検討"曽和美保子,
北野司, 近藤俊行, 森みどり, 鈴木敦志, 難波江宏一,上山智, 岩谷素顕, 竹内哲也,
赤﨑勇: 応用物理学会 SC 東海地区学術講演会 2013，名古屋大学，2013 年 11 月 16日.
138.
＊"トンネル接合を用いた多接合窒化物半導体の検討"黒川泰視, 合田智美, 加賀充, 岩谷素顕, 竹内哲也, 上山 智, 赤﨑 勇, 天野 浩: 電子情報通信学会 (ED,CPM, LQE 合同研究会）、大阪大学、2013 年 11 月 28 日-29 日．(*19 に対応)
139.
"350nm 紫外 LED 光取り出し効率改善に関する研究" 中嶋 翼,竹田健一郎, 岩谷
素顕, 上山 智, 竹内哲也, 赤﨑 勇, 天野 浩: 電子情報通信学会 (ED, CPM, LQE
合同研究会）、大阪大学、2013 年 11 月 28 日-29 日．
140.
＊"In situ X-ray diffraction monitoring of group III nitride growth by MOVPE"
Motoaki Iwaya, Hironori Kurokawa, Yosuke Katsu, Taiji Yamamoto, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki: Workshop on Ultra-Precision
Processing for III-Nitride, Santa Barbara, USA, October 16-18, 2013. [Invited Talk](*1 に対応)
141.
＊“Homoepitaxial growth of AlN films on freestanding AlN (0001) substrates bymetalorganic vapor phase epitaxy” Tomohiro Morishita, Motoaki Iwaya, Tetsuya
Takeuchi, Satoshi Kamiyama, Isamu Akasaki: SPIE Photonics West 2014, San
Francisco, USA, February 1-6, 2014. (*6 に対応)
142.
＊“Realization of the high conversion efficiency solar cells using high InN molar
fraction GaInN active layer” Motoaki Iwaya, Hironori Kurokawa, Yosuke Katsu, Taiji
Yamamoto, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki, Hiroshi Amano:
SPIE Photonics West 2014, San Francisco, USA, February 1-6, 2014. [InvitedTalk] (*19 に対応)
143.
“Improvement of photosensitivity in heterostructure-field-effect-transistor typephotosensors with a p-GaInN gate” Mami Ishiguro, Kazuya Ikeda, Takuya Murase,
Yuma Yamamoto, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki: ISPlasma,
Meijo University, Japan, March 2-6, 2014.
144.
“Nitride-based hetero-field-effect-transistor-type photosensors with extremely
high photosensitivity and complete solar-blind” Takuya Murase, Mami Ishiguro,（様式 2）
法人番号プロジェクト番号

231029S1201019

Yuma Yamamoto, Tetsuya Takeuchi, Satoshi Kamiyama, Isamu Akasaki: ISPlasma,
Meijo University, Japan, March 2-6, 2014.
145.
＊“Analysis of GaInN/GaN supperlattices on GaN template during growth byin-situ X-ray diffraction” Taiji Yamamoto, Daisuke Iida, Yasunari Kondo, Mihoko
Sowa, Shinya Umeda, Hiroyuki Matsubara, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: ISPlasma, Meijo University, Japan, March
2-6, 2014. (*1 に対応)
146.
“Realization of m-plane GaInN thick films grown on submicron patternedgroove-GaN substrate” Keiya Kozaki, Shinichiro Kondo, Shinya Umeda, Motoaki
Iwaya, Satoshi Kamiyama, Tetsuya Takeuchi, and Isamu Akasaki: ISPlasma, Meijo
University, Japan, March 2-6, 2014.
147.
"High reflective and low resistive silver electrode on p-GaN" Shunsuke Kawai,
Daisuke Iida, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu
Akasaki: ISPlasma, Meijo University, Japan, March 2-6, 2014.
148.
＊“Current confinement by nitride-based buried tunnel junction” Y. Kuwano, M.
Ino, T. Morita, D. Minamikawa, T. Takeuchi, S. Kamiyama, M. Iwaya, and I. Akasaki:
ISPlasma, Meijo University, Japan, March 2-6, 2014. (*18 に対応)
149.
＊“Low resistive GaInN tunnel junctions with high InN mole fractions” D.
Minamikawa, Y. Kuwano, S. Kawai, T. Morita, T. Takeuchi, S. Kamiyama, M. Iwaya,
and I. Akasaki: ISPlasma, Meijo University, Japan, March 2-6, 2014. (*5 に対応)
150.
“Improvement of the light extraction efficiency in 350-nm-emission UV
light-emitting diodes by novel distributed bragg reflector p-type electrode” T.
Nakashima, K. Takeda, M. Iwaya, T. Takeuchi, S. Kamiyama, I. Akasaki, and H.
Amano: ISPlasma, Meijo University, Japan, March 2-6, 2014.
151.
“Technologies of nitride-based LEDs”, Satoshi Kamiyama, ISPlasma, Meijo
University, Japan, March 2-6, 2014 [Tutorial talk]152.
＊“Observation of group III nitride semiconductors by in situ X-ray diffraction
monitoring during MOVPE growth” Motoaki Iwaya, Taiji Yamamoto, Daisuke Iida,
Mihoko Sowa, Yasunari Kondo, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu
Akasaki: International Conference on Materials and Characterization Techniques,
VIT University, Vellore, India, 10-12 March, 2014. [Invited talk] (*1 に対応)
153.
＊”表面プラズモンを利用した緑色 LED の高効率化に向けて” 飯田 大輔、岩谷 素
顕、竹内 哲也、上山 智、赤﨑 勇: 日本学術振興会ワイドギャップ半導体光・電子デバ
イス第 162 委員会 第 88 回研究会，主婦会館，2014 年 3 月 8 日 [依頼講演] (*11 に対応)
154.
“p 型 GaN 上の Ag 電極を用いた高反射電極の検討” 河合俊介，飯田大輔，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日.
155.
“電子オーバーフロー直接観測を用いた青色 LED 電子ブロック層の検討” 林健人，
松井健城，安田俊輝，勝野翔太，竹内哲也，上山智，岩谷素顕，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日.
156.
＊“分極制御による紫外発光素子のホール注入の検討” 安田俊輝，林健人，竹田
健一郎，中嶋翼，竹内哲也，岩谷素顕，上山智，赤﨑勇，天野浩: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*6 に対応)
157.
＊“Ⅲ族窒化物半導体トンネル接合上の高効率 LED” 森田隆敏，井野匡貴，桑野侑香，渡邉雅大，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第 61 回応用物理学会春季学
術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*5 に対応)

（様式 2）
法人番号プロジェクト番号

231029S1201019

158.
＊“III 族窒化物半導体埋込みトンネル接合による電流狭窄構造” 桑野侑香，堀川航佑，森田隆敏，井野匡貴，竹内哲也，上山智，岩谷素顕，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*18
に対応)
159.
＊“GaInN/GaN ヘテロ接合における緩和過程の転位密度依存性” 石原耕史，近藤保成，松原大幸，飯田大輔，山本泰司，曽和美保子，岩谷素顕，竹内哲也，上山智，
赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*1 に対応)
160.
＊“窒化物半導体 HFET 型 UV 光センサーの火炎センサー応用” 山本雄磨，村瀬卓
弥，石黒真未，山田知明，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*12
に対応)
161.
“高温スパッタ法により堆積した AlN バッファ層に関する検討” 伊藤弘晃，飯田大輔，近藤保成，曽和美保子，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日.
162.
＊“MOVPE 法を用いた高品質 AlInN の高速成長” 池山和希，小塚祐吾，安田俊輝，竹内哲也，上山智，岩谷素顕，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*3 に対応)
163.
＊“高成長速度 AlInN による窒化物半導体多層膜反射鏡” 小塚祐吾，池山和希，安田俊輝，竹内哲也，上山智，岩谷素顕，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*3 および*15 に対応)
164.
＊“非極性 m 面 GaN 基板上に作製した GaInN 系太陽電池の特性” 小崎桂矢，黒川泰視，岩谷素顕，上山智，竹内哲也，赤﨑勇，天野浩: 第 61 回応用物理学会春季学
術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*13 に対応)
165.
“蛍光 SiC の窒化物系白色 LED への応用に向けた研究” 上山智，岩谷素顕，竹内哲也，赤崎勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパ
ス、2014 年 3 月 17 日-20 日.[招待講演]
166.
＊“高 InN モル分率 GaInN を用いたトンネル接合” 南川大智，桑野侑香，河合俊介，森田隆敏，岩谷素顕，竹内哲也，上山智，赤崎勇: 第 61 回応用物理学会春季学術
講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*5 に対応)
167.
＊“その場観察 X 線回折法を用いた GaInN/GaN 超格子構造の最適化” 村瀬卓弥，
石黒真未，山田知明，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*1 に対応)
168.
＊“高感度な窒化物半導体 HFET 型ソーラーブラインド紫外光センサー” 村瀬卓弥，
石黒真未，山田知明，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年 3 月 17 日-20 日. (*13 に対応)
169.
“窒化物系青色 LED における光取り出し効率改善のためのレーザ加工方法の検討”
花井駿，鈴木敦志，北野司，梅田慎也，加藤貴久，上山智，竹内哲也，岩谷素顕，赤﨑勇: 第 61 回応用物理学会春季学術講演会、青山学院大学相模原キャンパス、2014 年
3 月 17 日-20 日.
170.
＊“Investigation of Hole Injection in UV-LEDs Utilizing Polarization Effect” T.
Yasuda , K. Hayashi, T. Nakashima, T. Takeuchi, S. Kamiyama, M. Iwaya, I.
Akasaki, and H. Amano: Conference on LED and Its Industrial Application ’14,
Yokohama, Japan, 22-24 April, 2014. (*7 に対応)
171.
＊“Roles of Mg Doping in Nitride Semiconductor-Based Light Emitting Diodes

（様式 2）
法人番号プロジェクト番号

231029S1201019

with Two Active Regions” K. Matsui, T. Morita, T. Suzuki, T. Takeuchi, S. Kamiyama,
M. Iwaya, and I. Akasaki: Conference on LED and Its Industrial Application ’14,
Yokohama, Japan, 22-24 April, 2014. (*5 に対応)
172.
＊“Nitride-Based Light Emitting Diodes with Buried Tunnel Junctions” M. Ino, Y.
Kuwano, T. Morita, D. Minamikawa, T. Takeuchi, S. Kamiyama, M. Iwaya, and I.
Akasaki: Conference on LED and Its Industrial Application ’14, Yokohama, Japan,
22-24 April, 2014. (*5 および*17 に対応)
173.
“Study of High-Reflective Ag-Based Electrode on p-Type GaN” S. Kawai, D.
Iida, M. Iwaya, T. Takeuchi, S. Kamiyama, and I. Akasaki: Conference on LED and
Its Industrial Application ’14, Yokohama, Japan, 22-24 April, 2014.
174.
“Fabrication of Nitride-Based Blue LED with Eliminating Light-Absorptive Laser
Scribing Damages” S. Hanai, I. Nakatani, A. Suzuki, T. Kitano, D. Iida, T. Kato, M.
Iwaya, T. Takeuchi, S. Kamiyama, and I. Akasaki: Conference on LED and Its
Industrial Application ’14, Yokohama, Japan, 22-24 April, 2014.
175.
＊“An alternative hole injection: Nitride-based tunnel junctions” T. Takeuchi, D.
Minamikawa, Y. Kuwano, M. Watanabe, M. Iwaya, S. Kamiyama, I. Akasaki:
International Conference on Metamaterials and Nanophysics, Varadero, Cuba, 22
April to 1 May 2014. [Invited Talk] (*5 に対応)
176.
＊“Development of nitride-based photosensor” Motoaki Iwaya, Takuya Murase,
Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: International
Conference on Metamaterials and Nanophysics, Varadero, Cuba, 22 April to 1 May
2014. [Invited Talk] (*13 に対応)
177.
＊ “Solar blind and exteremely high photosensitivity AlGaN-based UVhetero-field-effect-transistor-type photosensors” Takuya Murase, Mami Ishiguro,Tomoaki Yamada, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, Iasmu
Akasaki: Compound Semiconductor Week 2014, Montpellier, France, May 11-15,
2014. (*13 に対応)
178.
＊“Fabrication of Multi-junction GaInN-Based solar cell” Hironori Kurokawa,
Mitsuru Kaga, Tomomi Goda, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki, Hiroshi Amano: Compound Semiconductor Week 2014,
Montpellier, France, May 11-15, 2014. (*19 に対応)
179.
＊
“Activation
EnergyofExtremely
Low-Resistivity
andHigh-Carrier-Concentration Si-Doped AI0.05Ga0.95N” Motoaki Iwaya, Daisuke Iida,
Toru Sugiyama, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: 5thInternational Symposium on Growth of III-Nitrides, Atlanta, USA, May 18-22, 2014.
(*4 に対応)
180.
“Dislocation Density Dependence of Modal Gain in AlGaN/AlN Multiquantum
Wells” Tomoaki Yamada, Kenichiro Takeda, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, Isamu Akasaki, and Hiroshi Amano: 5th International
Symposium on Growth of III-Nitrides, Atlanta, USA, May 18-22, 2014.
181.
“Development of GaN-based semiconductors of excellent quality and and theirp-n junction blue-light-emitting devices”, Isamu Akasaki, Japan-Sweden Science,
Technology & Innovation Sympoium 2014, Stockholm, Sweden, May 21, 2014.
[Invited talk]
182.
＊“GaInN/GaN ヘテロ接合における緩和過程の転位密度依存性” 石原耕史, 近藤保成, 松原大幸, 岩谷素顕, 上山 智, 竹内哲也, 赤﨑 勇: 電子情報通信学会(ED・CPM・SDM 研究会), 名古屋大学、2014 年 5 月 28 日-29 日. (*1 に対応)
183.
＊“GaNSb の Sb 取り込みと表面形態に関する検討” 小森大資, 笹島浩希, 鈴木智行, 竹内哲也, 上山 智, 岩谷素顕, 赤崎 勇: 電子情報通信学会(ED・CPM・SDM 研（様式 2）
法人番号プロジェクト番号

231029S1201019

究会), 名古屋大学、2014 年 5 月 28 日-29 日. (*2 に対応)
184.
＊“窒化物半導体 HFET 型紫外受光素子の火炎センサー応用”山本雄磨, 村瀬卓
弥, 石黒真未, 山田知明, 岩谷素顕, 上山 智, 竹内哲也, 赤﨑 勇: 電子情報通信学会(ED・CPM・SDM 研究会), 名古屋大学、2014 年 5 月 28 日-29 日. (*13 に対応)
185.
＊“窒化物半導体 LED におけるキャリア輸送への分極固定電荷の影響” 勝野翔太,
林 健人, 安田俊輝, 岩谷素顕, 竹内哲也, 上山 智, 赤﨑 勇, 天野 浩: 電子情報通信学会(ED・CPM・SDM 研究会), 名古屋大学、2014 年 5 月 28 日-29 日. (*7 に対応)
186.
“窒化物系 LED 作製のレーザー加工による光取り出し効率向上の検討” 花井 駿,
鈴木敦志, 北野 司, 飯田大輔, 加藤貴久, 竹内哲也, 岩谷素顕, 上山 智, 赤﨑 勇:
通信学会(ED・CPM・SDM 研究会), 名古屋大学、2014 年 5 月 28 日-29 日.
187.
“モスアイ光取り出し構造を有する青色 LED”、上山 智: 光電相互変換第 125 委員会・産業用 LED 応用研究会合同研究会、名城大学（名城サテライト）、2014 年 7 月 11
日 [招待講演]
188.
＊“Dislocation density dependence of the critical thickness in GaInN/GaN
heterostructure” Taiji Yamamoto, Daisuke Iida, Yasunari Kondo, Mihoko Sowa,
Hiroyuki Matsubara, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, andIsamu Akasaki: 17th International Conference on Metalorganic Vapor Phase
Epitaxy, Lausanne, Switzerland, 13-18 July 2014. [Invited Talk] (*1 に対応)
189.
“AlGaN/AlN 多重量子井戸の光学利得の転位密度依存性” 山田知明, 岩谷素顕,
竹内哲也, 上山智, 赤﨑勇, 天野浩: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日.
190.
＊“非極性 m 面 GaN 基板上 GaInN 系太陽電池の特性” 小﨑桂矢, 黒川泰視, 山本泰司, 岩谷素顕, 竹内哲也, 上山智, 天野浩, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*13 に対応)
191.
＊“r 面サファイア基板上における a 面 GaN 成長に関する結晶性向上の検討” 加藤貴久, 水野尚之，伊藤弘晃，飯田大輔，藤井高志，福田承生，上山 智，竹内 哲也，岩谷 素顕，赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月
25-26 日. (*7 に対応)
192.
＊“ナノインプリント技術による表面プラズモン緑色 LED の作製と評価” 飯田大輔,
河合俊介, 土屋貴義, 岩谷素顕, 竹内哲也, 上山智, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*11 に対応)
193.
“MOVPE 選択成長による GaN ナノコラム“ 水野尚之, 上山智, 竹内哲也, 岩谷 素顕, 赤﨑 勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日.
194.
“電子オーバーフロー直接観測を用いた青色 LED 電子ブロック層の検討“ 林健人,
安田俊輝，勝野翔太, 竹内哲也, 上山智, 岩谷素顕, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日.
195.
＊“高感度な窒化物半導体 AlGaN 系 HFET 型ソーラーブラインド紫外光センサー“ 村瀬卓弥, 石黒真未, 山田知明, 山本雄磨, 岩谷素顕, 竹内哲也, 上山智, 赤﨑
勇: : 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*13 に対応)
196.
“p 型 GaN に対する Ag 電極上酸化保護層の検討“ 河合俊介, 飯田大輔, 岩谷素顕, 竹内哲也, 上山智, 赤﨑勇: : 第 6 回窒化物半導体結晶成長講演会, 名城大学,
2014 年 7 月 25-26 日.
197.
＊“窒化物半導体 AlGaN を用いた紫外受光素子の火炎センサー応用“ 山本雄磨,
村瀬卓弥, 山田知明, 岩谷素顕, 竹内哲也, 上山智, 赤﨑勇: : 第 6 回窒化物半導体
結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日(*13 に対応).

（様式 2）
法人番号プロジェクト番号

231029S1201019198.
＊“窒化物半導体 LED における分極電荷の補償“ 勝野翔太, 林健人, 安田俊輝,
岩谷素顕, 竹内哲也, 上山智, 天野浩, 赤﨑勇 : 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*7 に対応)
199.
＊“二つの活性層を有する発光ダイオードにおける Mg の影響“ 松井健城, 竹内哲也, 上山智, 岩谷素顕, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学,
2014 年 7 月 25-26 日. (*16 に対応)
200.
＊“窒化物半導体 LED における正孔伝導に対する分極電荷の影響“ 安田俊輝, 勝
野翔太, 林健人, 竹内哲也, 岩谷素顕, 上山智, 赤﨑勇, 天野浩: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*7 に対応)
201.
＊“Ⅲ族窒化物半導体による埋め込みトンネル接合“ 井野 匡貴, 南川 大智，水野尚之，竹内 哲也，上山 智，岩谷 素顕，赤﨑 勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*5 に対応)
202.
＊“GaInN 系トンネル接合を有する LED“ 南川大智, 井野匡貴, 岩谷素顕, 竹内哲也, 上山智, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月
25-26 日. (*5 に対応)
203.
“青色面発光レーザ共振器の内部損失の検討“ 中島啓介, 渡邉雅大, 堀川航佑,
竹内哲也, 上山智, 岩谷素顕, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日.
204.
“窒化物半導体面発光レーザにおける電流狭窄構造の特性評価“ 堀川航佑, 中島啓介, 小塚祐吾, 池山和希, 竹内哲也, 上山智, 岩谷素顕, 赤崎勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日.
205.
＊“その場観察 X 線回折法を用いた GaInN 系太陽電池構造の最適化“ 山本泰司,
黒川泰視, 小﨑桂矢, 岩谷素顕, 上山智, 竹内哲也, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*1 に対応)
206.
＊“高成長速度 AlInN 用いた多層膜反射鏡に関する研究“ 小塚祐吾, 池山和希,
安田俊輝, 竹内哲也, 上山智, 岩谷素顕, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*3 および*15 に対応)
207.
＊“GaInN/GaN ヘテロ接合における緩和過程の転位密度依存性“ 石原耕史, 近藤保成, 松原大幸, 岩谷素顕, 上山智, 竹内哲也, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*1 に対応)
208.
＊“高品質 AlInN の高速成長の検討“ 池山和希, 小塚祐吾, 安田俊輝, 竹内哲也,
上山智, 岩谷素顕, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014
年 7 月 25-26 日. (*3 に対応)
209.
＊“MOCVD 法を用いた GaNSb の結晶成長“ 小森大資, 笹島浩希, 鈴木智行, 竹
内哲也, 岩谷素顕, 上山智, 赤﨑勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*2 に対応)
210.
＊“MOCVD 法による Sb を添加した AlN および GaN の作製“ 笹島 浩希, 小森 大資, 竹内 哲也, 岩谷 素顕, 上山 智, 赤﨑 勇: 第 6 回窒化物半導体結晶成長講演会, 名城大学, 2014 年 7 月 25-26 日. (*2 に対応)
211.
“窒化物系 LED 技術最前線” 上山智: 第 6 回窒化物半導体結晶成長講演会, 名城
大学, 2014 年 7 月 25-26 日. [チュートリアル講演]
212.
＊“Lattice relaxation process in GaInN/GaN heterostructure system as functionof dislocation density in underlying GaN layer” Motoaki Iwaya, Koji Ishihara, Taiji
Yamamoto, Daisuke Iida, Yasunari Kondo, Hiroyuki Matsubara, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: The Workshop on Ultra-Precision
Processing for Wide-Gap Semiconductors 2014, Bath, United Kingdom, August（様式 2）
法人番号プロジェクト番号

231029S1201019

20-22 2014. [Invited talk] (*1 に対応)
213.
＊ “Nitride-based optoelectronic devices utilizing tunnel junctions” Tetsuya
Takeuchi, Motoaki Iwaya, Satoshi Kamiyama, and Isamu Akasaki: Lester Eastman
Conference on High Performance Devices 2014, Ithaca, USA August 5-7 2014.
[Invited Talk] (*5 および*16 に対応)
214.
＊“Demonstration of GaInN-based laser pumped by an electron beam” M.
Iwaya, K. Kozaki, T. Yamada, T. Takeuchi, S. Kamiyama, I. Akasaki, Y. Honda, H.
Amano, S. Iwayama, J. Matsuda, N. Matsubara, T. Matsumoto: International
Workshop on Nitride semiconductor 2014, Wrocław, Poland August 24-29 2014.
[Invited Talk] (*14 に対応)
215.
＊ “Nitride-based tunnel junctions as an alternative hole injection” Tetsuya
Takeuchi, Motoaki Iwaya, Satoshi Kamiyama, and Isamu Akasaki: International
Workshop on Nitride semiconductor 2014, Wrocław, Poland August 24-29 2014.
[Invited Talk] (*5 および*16 に対応)
216.
＊“Laser lift-off technique of 380-nm light emitting diodes grown on GaN
substrate” D. Iida , S. Kawai , T. Tsuchiya, N. Ema, M. Iwaya, T. Takeuchi, S.
Kamiyama, and I. Akasaki: International Workshop on Nitride semiconductor 2014,
Wrocław, Poland August 24-29 2014. (*18 に対応)
217.
＊ “Fabrication of surface plasmon based green light emitting diodes bynanoimprint technique” D. Iida , S. Kawai, T. Tsuchiya, M. Iwaya, T. Takeuchi, S.
Kamiyama, and I. Akasaki: International Workshop on Nitride semiconductor 2014,
Wrocław, Poland August 24-29 2014. (*11 に対応)
218.
＊“Fabrication of nonpolar m-plane GaInN-based solar cell on freestanding
m-plane GaN substrate” K. Kozaki, H. Kurokawa, M. Iwaya, T. Takeuchi, S.
Kamiyama, and I. Akasaki: International Workshop on Nitride semiconductor 2014,
Wrocław, Poland August 24-29 2014. (*5 に対応)
219. “Lattice relaxation process in GaInN/GaN heterostructure system as function of
dislocation density in underlying GaN layer” Motoaki Iwaya, Koji Ishihara, Taiji
Yamamoto, Daisuke Iida, Yasunari Kondo, Hiroyuki Matsubara, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: Workshop on Ultra-Precision Processing
for Wide Band Gap Semiconductors, Bath, UK, 20-22 Augast, 2014.
220. “Low Resistive and Low Absorptive Nitride-Based Tunnel Junctions” Daichi
Minamikawa, Masataka Ino, Daiki Takasuka, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: Material Research Society Fall Meeting,
Boston, USA, Nov. 30-Dec. 5, 2014.
221. “Evaluation of Electron Overflow in Blue-LEDs with P-AlGaN Or P-GaN Electron
Blocking Layer“ Kento Hayashi, Toshiki Yasuda, Shota Katsuno, Tetsuya Takeuchi,
Satoshi Kamiyama, Motoaki Iwaya, Isamu Akasaki: Material Research Society
Fall Meeting, Boston, USA, Nov. 30-Dec. 5, 2014.
222. “Optimization of Carrier Distributions in Periodic Gain Structures towards BlueVCSELs” Kenjo Matsui, Kosuke Horikawa, Yugo Kozuka, Kazuki Ikeyama, Tetsuya
Takeuchi, Satoshi Kamiyama, Motoaki Iwaya, Isamu Akasaki: Material ResearchSociety Fall Meeting, Boston, USA, Nov. 30-Dec. 5, 2014.
223. “Study on the High Reflective Ag Ohmic Contact on High Carrier Concentration
Si-Doped Al0.03Ga0.97N” Shunsuke Kawai, Daisuke Iida, Motoaki Iwaya, Tetsuya

（様式 2）
法人番号プロジェクト番号

231029S1201019

Takeuchi, Satoshi Kamiyama, Isamu Akasaki: Material Research Society Fall
Meeting, Boston, USA, Nov. 30-Dec. 5, 2014.
224. “Analysis of GaInN/GaN Heterostructure Grown on GaN Underlying Layer withDifferent Dislocation Density” Koji Ishihara, Taiji Yamamoto, Daisuke Iida, Yasunari
Kondo, Hiroyuki Matsubara, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama,
Isamu Akasaki: Material Research Society Fall Meeting, Boston, USA, Nov.
30-Dec. 5, 2014.
225. “Growths of AlInN Single Layers and Distributed Bragg Reflectors for VCSELs”
Yugo Kozuka, Kazuki Ikeyama, Toshiki Yasuda, Tetsuya Takeuchi, Satoshi
Kamiyama, Motoaki Iwaya, Isamu Akasaki: Material Research Society Fall Meeting,
Boston, USA, Nov. 30-Dec. 5, 2014.
226. “Laser lift-off technique for freestanding GaNsubstrate using an In droplet formedby thermaldecomposition of GaInN and its application to LED” Motoaki Iwaya,
Daisuke Iida, Syunsuke Kawai, Nobuaki Ema, Takayoshi Tsuchiya, Tetsuya
Takeuchi, Satoshi Kamiyama, and Isamu Akasaki: 10th International Symposium
on Semiconductor Light Emitting Devices (ISSLED 2014) Kaohsiung, Taiwan, Dec.
14th-19th, 2014.
227. “In-situ x-ray diffraction analysis for MOVPE growth of nitride semiconductors”
Motoaki Iwaya, Taiji Yamamoto, Koji Ishihara, Tetsuya Takeuchi, Satoshi
Kamiyama, and Isamu Akasaki: SPIE Photonics West, San Fransisco, USA, Feb.
7th-12th, 2015.
228. “Performance of nitride-based blue LED fabricated on sapphire substrate withnanostructured SiO2” Shun Hanai, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki, and Tsukasa Kitano: SPIE Photonics West, San
Fransisco, USA, Feb. 7th-12th, 2015.
229. “Analysis of GaInN/GaN Superlattice on GaN by in situ X-Ray Diffraction
Monitoring Attached with Metalorganic Vapor Phase Epitaxy Equipment” Motoaki
Iwaya, Taiji Yamamoto, Daisuke Iida, Koji Ishihara, Tetsuya Takeuchi, Satoshi
Kamiyama, Isamu Akasaki: Material Research Society Spring Meeting, San
Francisco, USA, April 6-10, 2015.
230. “Electrical Properties of High CarrierConcentration n-Type AlGaN” K. Takeda, K.
Mori, T. Kusafuka, M. Iwaya, T. Takeuchi, S. Kamiyama, and I. Akasaki: The 3rdInternational Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
231. “Correlation between Crystal Qualities and Electrical Properties in Si-Doped
Al0.6Ga0.4N” T. Yasuda, S. Katsuno, T. Takeuchi, M. Iwaya, S. Kamiyama, I. Akasaki,
and H. Amano: The 3rd International Conference on Light-Emitting Devices andTheir Industrial Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th,2015.
232. “AlN Epitaxial Growth on Sapphire with an Intermediate Layer” S. Katsuno, T.
Yasude, M. Iwaya, T. Takeuchi, S. Kamiyama, I. Akasaki, and H. Amano: The 3rdInternational Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.

（様式 2）
法人番号プロジェクト番号

231029S1201019

233. “AlN Growth on Sputtering AlN Template Substrate by Hydride Vapor Phase
Epitaxy” D. Yasui, H. Miyake, K. Hiramatsu, M. Iwaya, I. Akasaki, and H. Amano:
The 3rd International Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
234. “MOVPE Growth of AlNSb Alloys” K. Suzuki, D. Komori, H. Sasajima, K. Takarabe,T. Takeuchi, M. Iwaya, S. Kamiyama, and I. Akasaki: The 3rd International
Conference on Light-Emitting Devices and Their Industrial Applications (LEDIA '15),
Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
235. “Carrier Gas Dependence on GaNSb MOVPE Growth” D. Komori, H. Sasajima, K.
Takarabe, K. Suzuki, T. Takeuchi, S. Kamiyama, M. Iwaya, and I. Akasaki: The 3rdInternational Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
236. “Electrical Properties of GaNSb Grown at Low Temperatures” K. Takarabe, D.
Komori, K. Suzuki, T. Takeuchi, M. Iwaya, S. Kamiyama, and I. Akasaki: The 3rdInternational Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
237. “Contact Characteristics of V-Based Electrode for High AlN Molar Fraction
n-AlGaN” K. Mori, K. Takeda, T. Kusafuka, M. Iwaya, T. Takeuchi, S. Kamiyama, I.
Akasaki, and H. Amano: The 3rd International Conference on Light-Emitting
Devices and Their Industrial Applications (LEDIA '15), Pacifico Yokohama, Japan,
Apr. 22nd-24th, 2015.
238. “Nitride-Based Tunnel Junctions towards Deep UV-LEDs” D. Takasuka, D.
Minamikawa, M. Ino, T. Takeuchi, M. Iwaya, S. Kamiyama, H. Amano, and I.
Akasaki: The 3rd International Conference on Light-Emitting Devices and Their
Industrial Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
239. “Optimization of Periodic Gain Structures toward Blue VCSELs” K. Matsui, D.
Komori, K. Horikawa, T. Takeuchi, S. Kamiyama, M. Iwaya, and I. Akasaki: The 3rdInternational Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
240. “Conductive n-Type AlInN/GaN Distributed Bragg Reflectors” K. Ikeyama, Y.
Kozuka, T. Yasuda, T. Takeuchi, S. Kamiyama, M. Iwaya, and I. Akasaki: The 3rdInternational Conference on Light-Emitting Devices and Their Industrial
Applications (LEDIA '15), Pacifico Yokohama, Japan, Apr. 22nd-24th, 2015.
241. “A Pulsed Operation of VCSELs on AlInN/GaN DBR Grown with High GrowthRate” Y. Kozuka, K. Ikeyama, T. Akagi, S. Iwayama, K. Nakata, T. Takeuchi, S.
Kamiyama, M. Iwaya, I. Akasaki: The 3rd International Conference on
Light-Emitting Devices and Their Industrial Applications (LEDIA '15), Pacifico
Yokohama, Japan, Apr. 22nd-24th, 2015.
242. “Electrical properties of extremely low-resistivity and high-carrier-concentration
Si-doped AlGaN with low AlN molar fraction” Kunihiro Takeda, Kazuki Mori, Toshiki
Kusafuka, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi Kamiyama, and Isamu
Akasaki: The 7th Asia-Pacific Workshop on Widegap Semiconductors (APWS
2015), Seoul, Korea, May 17-20, 2015.

（様式 2）
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243. “Growth of AlN layer on sputtered AlN template substrate by hydride vapor phase
epitaxy” Shinya Tamaki, Daiki Yasui, Hideto Miyake, Kazumasa Hiramatsu, Motoaki
Iwaya, Isamu Akasakai, and Hiroshi Amano: The 7th Asia-Pacific Workshop on
Widegap Semiconductors (APWS 2015), Seoul, Korea, May 17-20, 2015.
244. “In situ X-ray diffraction analysis of GaInN/GaN heterostructure during MOVPE
growth” Motoaki Iwaya, Taiji Kondo, Koji Ishihara, Junya Osumi, Tetsuya Takeuchi,
Satoshi Kamiyama, and Isamu Akasaki: The 7th Asia-Pacific Workshop on
Widegap Semiconductors (APWS 2015), Seoul, Korea, May 17-20, 2015.
245. “Extremely Low-resistivity and High-carrierconcentration Si-doped AlGaN with
Low AlN Molar Fraction for Improvement of Wall Plug Efficiency of Nitride-basedLED” Motoaki Iwaya, Kunihiro Takeda, Daisuke Iida, Tetsuya Takeuchi, Satoshi
Kamiyama, and Isamu Akasaki: The 11th Conference on Lasers and Electro-Optics
Pacific Rim (CLEO-PR 2015), Busan, Korea, August 24-28, 2015.
246. “Extremely Low-resistivity and High-carrierconcentration Si-doped AlGaN with
Low AlN Molar Fraction for Improvement of Wall Plug Efficiency of Nitride-basedLED” Motoaki Iwaya, Kunihiro Takeda, Daisuke Iida, Tetsuya Takeuchi, Satoshi
Kamiyama, and Isamu Akasaki:The 11th International Conference on Nitride
Semiconductor (ICNS-11), Beijiin, China, Augast 30-September 4, 2015.
247. “MOVPE法を用いたSi添加AlInN/GaNヘテロ接合の電気的特性” 池山和希，安田俊輝，小塚祐吾，堀川航佑，竹内哲也，上山智，岩谷素顕，赤﨑勇: 第75回 応用物
理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
248.
“MgドープAlGaNによるAlGaN/GaN HFET型紫外フォトセンサーの高性能化”
山本雄磨，村瀬卓弥，草深敏匡，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
249.
“SiCの陽極酸化プロセスにおける光照射に関する検討” 日高公崇香，渡和樹，
岩谷素顕，上山智，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，
2014年9月17日-9月20日.
250.
“GaInN/GaNヘテロ接合における格子緩和の評価” 石原耕史，近藤保成，松原大幸，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
251.
“高InNモル分率GaInNを用いたトンネル接合（2）”南川大智，井野匡貴，高須賀大貴，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
252.
“MOVPE選択成長によるGaNナノコラムのマスクパターン依存性” 水野尚之，
加藤嵩裕，上山智，竹内哲也，岩谷素顕，赤崎勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
253.
“高キャリア濃度Siドープn型Al0.03Ga0.97N上のAg電極を用いた高反射電極の検討” 河合俊介，飯田大輔，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
254.
“III族窒化物半導体トンネル接合を用いた電流狭窄構造の低抵抗化” 井野匡貴，南川大智，竹内哲也，上山智，岩谷素顕，赤﨑勇: 第75回 応用物理学会秋季
学術講演会，北海道大学，2014年9月17日-9月20日.
255.
“レーザーリフトオフ技術によるGaN 基板を剥離した近紫外LED” 飯田大輔，

（様式 2）
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河合俊介，土屋貴義，江間伸明，岩谷素顕，竹内哲也，上山智，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
256.
“GaInN/GaN超格子を用いた窒化物太陽電池構造の検討” 山本泰司，黒川泰視，
小﨑桂矢，岩谷素顕，上山智，竹内哲也，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
257.
“r面サファイア基板上におけるa面GaN成長に関する検討” 加藤貴久，水野尚之，伊藤弘晃，飯田大輔，藤井高志，福田承生，上山智，竹内哲也，岩谷素顕，赤﨑勇: 第75回 応用物理学会秋季学術講演会，北海道大学，2014年9月17日-9月20日.
258.
“AlInNエピタキシャル層の高速成長と多層膜反射鏡への応用” 池山和希、竹内哲也、岩谷素顕、上山智、赤﨑勇：第44回日本結晶成長学会国内会議、学習院大学、2014年11月6日-8日.
259.
“スパッタ法AlNテンプレート基板を用いたAlNのHVPE成長” 安井 大貴、三宅秀人、平松 和政、岩谷 素顕、赤崎 勇、天野 浩: 第62回応用物理学会春季学術講
演会、東海大学湘南キャンパス、2015年3月11日-14日.
260.
“高キャリア濃度n型AlGaNの電気的特性” 武田 邦宏、森 一喜、山田 知明、
岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇: 第62回応用物理学会春季学術講演会、
東海大学湘南キャンパス、2015年3月11日-14日.
261.
“深紫外LEDに向けた窒化物半導体トンネル接合の検討” 高須賀 大貴、南川大智、井野 匡貴、竹内 哲也、岩谷 素顕、上山 智、天野 浩、赤﨑 勇: 第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
262.
“トンネル接合を用いたモノリシック型白色LEDの検討” 加藤 貴久、飯田 大輔、河合 俊介、上山 智、竹内 哲也、岩谷 素顕、赤﨑 勇: 第62回応用物理学会
春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
263.
“GaN上に作製したGaInN単膜およびGaInN/GaN超格子のX線その場観察測定” 大角 純也、石原 耕史、山本 泰司、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇:
第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
264.
“表面プラズモン共鳴のナノ構造と屈折率の依存性の検討” 丹羽 一将、加藤
貴久、河合 俊介、飯田 大輔、上山 智、竹内 哲也、岩谷 素顕、赤崎 勇: 第62回
応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
265.
“高Al組成のn-AlGaNにおけるV系電極のコンタクト特性” 森 一喜、武田 邦
宏、草深 敏匡、岩谷 素顕、上山 智、竹内 哲也、赤崎 勇、天野 浩: 第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
266.
“高性能近紫外LEDの作製に関する検討” 弦間 彩花、河合 俊介、岩谷 素顕、
竹内 哲也、上山 智、赤﨑 勇: 第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
267.
“SiO2ナノ周期構造を有する基板上に作製した窒化物系青色LEDの発光特性”
花井 駿、北野 司、岩谷 素顕、竹内 哲也、上山 智、赤崎 勇: 第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
268.
“ナノ構造導入によるLEDの性能向上に関する研究” 上山 智、岩谷 素顕、竹内 哲也、赤﨑 勇: 第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.

（様式 2）
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269.
“高成長速度AlInN/GaN多層膜反射鏡を用いた面発光レーザのパルス発振” 小塚 祐吾、池山 和希、赤木 孝信、岩山 章、中田 圭祐、竹内 哲也、上山 智、岩谷 素顕、赤﨑 勇: 第62回応用物理学会春季学術講演会、東海大学湘南キャンパス、2015年3月11日-14日.
270.
“青色面発光レーザに向けたGaInN 活性層に関する研究” 古田貴士, 松井健
城，堀川航佑，小塚祐吾，池山和樹，岩山 章，赤木孝信，竹内哲也，上山 智，岩谷素顕，赤﨑 勇: 第7回窒化物半導体結晶成長講演会、東北大学 片平キャンパス、2015年5月7-8日.
271.
“AlInN 膜およびAlInN/GaN 多層膜反射鏡の電気的特性” 吉田翔太朗，池山和希，竹内哲也，上山 智，岩谷素顕，赤﨑 勇: 第7回窒化物半導体結晶成長講演会、
東北大学 片平キャンパス、2015年5月7-8日.
272.
“その場観察X 線回折法を用いた GaN 上のGaInN 単膜およびGaInN/GaN
超格子の評価” 大角純也，石原耕史，山本泰司，岩谷素顕，竹内哲也，上山 智，
赤﨑 勇: 第7回窒化物半導体結晶成長講演会、東北大学 片平キャンパス、2015
年5月7-8日.
273.
“超高キャリア密度n 型AlGaN の結晶成長と電気伝導特性” 岩谷素顕，武田邦宏，飯田大輔，杉山 徹，竹内哲也，上山 智，赤﨑 勇: 第7回窒化物半導体結晶成長講演会、東北大学 片平キャンパス、2015年5月7-8日.
274.
“窒化物半導体面発光レーザの現状と照明応用に向けた将来展望” 平成 27 年度（第 48 回）照明学会全国⼤会 固体光源分科会シンポジウム 2015 年 8 月 27 日.
275.
T. Takeuchi, M. Iwaya, S. Kamiyama, and I. Akasaki, "GaInN-based tunnel
junctions for novel optoelectronic devices", EMN Meeting on Photovoltaics, 2016,
Hong Kong.
276.
T. Takeuchi, M. Iwaya, S. Kamiyama, and I. Akasaki, "Current status of
GaN-based vertical-cavity surface-emitting lasers (VCSELs)", 4th International
Workshop on LEDs and Solar Applications, 2016, Nagoya.
277.
T. Yasuda, S. Yoshida, T. Takeuchi, M. Iwaya, S. Kamiyama, I. Akasaki, and H.
Amano, Polarization Induced Hole Accumulations in Nitride Semiconductor
Heterostructures, LEDIA'16, Yokohama, Japan, (2016)
278.
K. Matsui, K. Ikeyama, T. Furuta, Y. Kozuka, T. Akagi, T. Takeuchi, S.
Kamiyama, M. Iwaya, and I. Akasaki, GaN-based VCSELs using Periodic Gain
Structures, LEDIA'16, Yokohama, Japan, (2016)
279.
N. Kuwabara, T. Yasuda, S. Katsuno, N. Koide, T. Takeuchi, M. Iwaya, S.
Kamiyama, and I. Akasaki, ITO/Ga2O3 Multilayer Electrodes Towards Deep
UV-LEDs , LEDIA'16, Yokohama, Japan, (2016)
280.
S. Ushida, A. Yoshikawa, M. Iwaya, T. Takeuchi, S. Kamiyama, and I. Akasaki,
Temperature Dependence of the Nitride-based HFET Structure Photosensors,
LEDIA'16, Yokohama, Japan, (2016)
281.
K. Suzuki, K. Takarabe, D. Komori, D. Takasuka, N. Koide, T. Takeuchi, M.
Iwaya, S. Kamiyama, and I. Akasaki, Low-Temperature Grown p-Side Structure withGaInN Tunnel Junction and n-GaNSb, LEDIA'16, Yokohama, Japan, (2016)
282.
T. Senga, N. Nagata, M. Iwaya, T. Takeuchi, S. Kamiyama, I. Akasaki, Si（様式 2）
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concentration dependence of laser oscillation characteristics in AlGaN multiple
quantum well active layer, ICCGE-18, Nagoya, Japan (2016).
283.
J. Osumi, R. Kanayama, M. Iwaya, T. Takeuchi, S. Kamiyama, I. Akasaki,
Observation of AlGaN/GaN heterostructure by in situ XRD attached metalorganic
vapor phase epitaxial equipment, ICCGE-18, Nagoya, Japan (2016).
284.
Y. Akatsuka, D. Takasuka, M. Ino, T. Akagi, T. Takeuchi, M. Iwaya, S.
Kamiyama, I. Akasaki, Low resistiveGaInN tunneljunctionswith high Si
concentrations, ICCGE-18, Nagoya, Japan (2016).
285.
Kenjo Matsui, Takashi Furuta, Natsumi Hayashi, Yugo Kozuka, Takanobu
Akagi, Tetsuya Takeuchi, Satoshi Kamiyama, Motoaki Iwaya and Isamu Akasaki,
3-mW RT-CW GaN-Based VCSELs and Their Temperature Dependence,
IWN2016, Florida, USA (2016).
286.
Daiki Jinno, Shun Otsuki, Teruyuki Niimi, Hisayoshi Daicho, Motoaki Iwaya,
Tetsuya Takeuchi, Satoshi Kamiyama and Isamu Akasaki, High Temperature
Annealing of Sputtered AlN Buffer Layer on r-Plane Sapphire Substrate and its
Effect on Crystalline Quality of a-Plane GaN, IWN2016, Florida, USA (2016).
287.
Bo Monemar, Plamen P. Paskov, K. Takeda, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama and Isamu Akasaki, Low Temperature Photoluminescence inHighly Si-Doped AlxGa1-xN with x < 0.09, IWN2016, Florida, USA (2016).
288.
Yuki Kurisaki, Satoshi Kamiyama, Tetsuya Takeuchi, Motoaki Iwaya and Isamu
Akasaki, Theoretical Investigation of Nitride Nanowire-Based QuantumShell
Lasers, IWN2016, Florida, USA (2016).
289.
Ryoma Seiki, Daisuke Komori, Kazuki Ikeyama, Toshiaki Ina, Takeyoshi
Onuma, Takao Miyajima, Tetsuya Takeuchi, Satoshi Kamiyama, Motoaki Iwaya andIsamu Akasaki, Local Structural Analysis around In Atoms in Al0.82In0.18N alloy by
Using X-Ray Absorption Fine-Structure Measurements, IWN2016, Florida, USA(2016).
290.
Ryousuke Kanayama, Junya Osumi, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama and Isamu Akasaki, Characterization of AlGaN/GaN Heterostructure by
In Situ X-Ray Diffraction Attached Metal Organic Vapor Phase Epitaxy, IWN2016,
Florida, USA (2016).
291.
Takao Miyajima, Daisuke Komori, Toshiaki Ina, Ryoma Seiki, Kiyofumi Nitta,
Tetsuya Takeuchi, Tomoya Uruga, Satoshi Kamiyama, Motoaki Iwaya and Isamu
Akasaki, Determination of the Site of Sb Occupation in MOCVD-Grown GaN1xSbx
Using X-Ray Absorption Fine-Structure Measurements, IWN2016, Florida, USA(2016).
292.
Yasuto Akatsuka, Daiki Takasuka, Takanobu Akagi, Tetsuya Takeuchi, Motoaki
Iwaya, Satoshi Kamiyama and Isamu Akasaki, Buried Tunnel Junctions Using Low
Resistive GaInN Tunnel Junctions with High Si Concentrations, IWN2016, Florida,
USA (2016).
293.
Akira Yoshikawa, Saki Ushida, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama and Isamu Akasaki, Physical Property of the High Photosensitive Field
Effect Transistor Type UV Photosensors with AlGaN/AlGaN Hetero Structure,（様式 2）
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IWN2016, Florida, USA (2016).
294.
Saki Ushida, Akira Yoshikawa, Motoaki Iwaya, Satoshi Kamiyama, Tetsuya
Takeuchi and Isamu Akasaki, Temperature Dependence in AlGaN-Based
Heterostructure Field-Effect Transistor Type UV Photosensors, IWN2016, Florida,
USA (2016).
295.
Akira Yoshikawa, Saki Ushida, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama and Isamu Akasaki, Physical Property of the High Photosensitive Field
Effect Transistor Type UV Photosensors with AlGaN/AlGaN Hetero Structure,
IWN2016, Florida, USA (2016).
296.
Junichiro Ogimoto, Yugo Kozuka, Takanobu Akagi, Natsumi Hayashi, Tetsuya
Takeuchi, Satoshi Kamiyama, Motoaki Iwaya and Isamu Akasaki, Design andFabrication of Modulation-Doped GaN-Based Vertical Cavities for Blue
Surface-Emitting Lasers, IWN2016, Florida, USA (2016).
297.
Toshiki Yasuda, Tetsuya Takeuchi, Motoaki Iwaya, Satoshi Kamiyama, Isamu
Akasaki, Hole Accumulations to Polarization Charges in Relaxed AlGaN
Heterostructures with High AlN Mole Fractions, IWN2016, Florida, USA (2016).
298.
Noriaki Nagata1, Takashi Senga, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama and Isamu Akasaki, Low Resistivity Ohmic Contact V-Based ElectrodeContributed by Using Thin SiNx Intermediate Layer for High AlN Molar Fraction
n-Type AlGaN, IWN2016, Florida, USA (2016).
299.
新實 輝行、神野 大樹、岩谷 素顕、竹内 哲也、上山 智、赤崎 勇、r 面サファイア基板上 a 面 GaInN 量子井戸構造に関する検討、第 63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
300.
栗崎 湧気、上山 智、竹内 哲也、岩谷 素顕、赤﨑 勇、Ⅲ族窒化物ナノワイ
ヤを用いた端面発光レーザの理論的検討、第 63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
301.
千賀 崇史、永田 訓章、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇、AlGaN 多重量子井戸構造のレーザ発振特性の Si 濃度依存性、第 63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
302.
高西 徹、大角 純也、岩谷 素顕、上山 智、竹内 哲也、赤﨑 勇、半極性(10-11)
面自立 GaN 基板上 GaInN の成長機構の解析、第 63 回応用物理学会春季学術講演
会シンポジウム、東京、2015.
303.
岩瀬 航平、Lim Stanley、澁谷 弘樹、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇、Core-Shell 型 MQW における発光波長の成長条件依存性、第 63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
304.
袴田 淳哉、草深 敏匡、千賀 崇史、岩谷 素顕、竹内 哲也、上山 智、三宅 秀
人、赤崎 勇、アニール処理 AlN 下地層上 AlGaN/AlN-MQW の光学特性、第 63 回
応用物理学会春季学術講演会シンポジウム、東京、2015.
305.
大角 純也、金山 亮介、髙西 徹、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇、
その場観察技術を用いた MOVPE 法で作製する AlGaN/GaN ヘテロ構造の観察、第63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
306.
牛田 彩希、吉川 陽、山本 雄磨、奥村 俊紀、岩谷 素顕、上山 智、竹内 哲
也、赤﨑 勇、窒化物半導体 HFET 型光センサの温度特性、第 63 回応用物理学会春（様式 2）
法人番号プロジェクト番号

231029S1201019

季学術講演会シンポジウム、東京、2015.
307.
永田 訓章、森 一喜、武田 邦宏、草深 敏匤、岩谷 素顕、竹内 哲也、上山 智、
赤﨑 勇、Si3N4 による高 AlN モル分率 n-AlGaN 用 V 系電極の低接触比抵抗化、第63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
308.
池山 和希、小塚 祐吾、吉田 翔太朗、松井 健城、赤木 孝信、岩山 章、竹内
哲也、上山 智、岩谷 素顕、赤﨑 勇、n 型 AlInN/GaN DBR を有する微小共振器
LED、第 63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
309.
鈴木 健太、財部 覚、小森 大資、高須賀 大貴、小出 典克、竹内 哲也、岩谷素顕、上山 智、赤﨑 勇、GaInN トンネル接合と n 型 GaNSb による低温 p 側構造の作製、第 63 回応用物理学会春季学術講演会シンポジウム、東京、2015.
310.
桑原 奈津子、安田 俊輝、勝野 翔太、小出 典克、竹内 哲也、岩谷 素顕、上山 智、赤﨑 勇、深紫外 LED のための ITO/Ga2O3 多層膜透明電極の検討、第 63回応用物理学会春季学術講演会シンポジウム、東京、2015.
311.
K. Matsui, T. Furuta, Y. Kozuka, T. Akagi, T. Takeuchi, S. Kamiyama, M. Iwaya
and I. Akasaki, GaN-based VCSEL using a periodic gain structure consisting of two
GaInN 5QWs, 35th Electronic Materilas Symposium, Biwako, Japan (2016).
312.
T. Yasuda, S. Katsuno, N. Kuwabara, T. Takeuchi, M. Iwaya, S. Kamiyama, I.
Akasaki, Improvement of p-type electrical property by polarization-doping in
graded-AlGaN layer, 35th Electronic Materilas Symposium, Biwako, Japan (2016).
313.
袴田 淳哉、岩谷 素顕、上山 智、竹内 哲也、三宅 秀人、赤﨑 勇, 高温アニール処理のスパッタ AlN 層における効果、 第 8 回 窒化物半導体結晶成長講演会、
京都（2016）314.
鈴木 健太、財部 覚、高須賀 大貴、小出 典克、 竹内 哲也、岩谷 素顕、上山 智、赤﨑 勇、GaInN トンネル接合と n 型 GaNSb による低温 p 側構造の作製、第 8 回 窒化物半導体結晶成長講演会、京都（2016）315.
吉川 陽、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇、 Nano-size concavo-convex(NCC)サファイア基板を用いた高品質 AlN 膜成長、第 8 回 窒化物半導体結晶成長
講演会、京都（2016）316.
牛田 彩希、吉川 陽、岩谷 素顕、上山 智、竹内 哲也、赤﨑 勇、窒化物半導体 HFET 型光センサの動作に関する温度依存性、第 8 回 窒化物半導体結晶成長講
演会、京都（2016）317.
鈴木 健太、財部 覚、高須賀 大貴、小出 典克、 竹内 哲也、岩谷 素顕、上山 智、赤﨑 勇、GaInN トンネル接合と n 型 GaNSb による低温 p 側構造の作製、第 8 回 窒化物半導体結晶成長講演会、京都（2016）318.
上山 智、竹内 哲也、岩谷 素顕、赤﨑 勇、窒化物系ナノワイヤーおよび量子殻構造の作製と、光デバイス応用、第 76 回応用物理学会秋季学術講演会シンポジ
ウム、名古屋、2015319.
牛田 彩希、吉川 陽、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇、AlGaN 系
HFET 型光センサの SiNx パッシベーション効果、第 76 回応用物理学会秋季学術講
演会シンポジウム、名古屋、2015320.
安田 俊輝、竹内 哲也、岩谷 素顕、上山 智、赤﨑 勇、天野 浩、分極電荷による高ホール濃度を有する p 型 AlGaN の作製、第 76 回応用物理学会秋季学術講演
会シンポジウム、名古屋、2015

（様式 2）
法人番号プロジェクト番号

231029S1201019321.
澁谷 弘樹、上山 智、岩谷 素顕、竹内 哲也、赤﨑 勇、低温下におけるⅢ族
窒化物ナノワイヤの埋め込み成長に関する研究 、第 76 回応用物理学会秋季学術講
演会シンポジウム、名古屋、2015322.
金山 亮介、大角 純也、岩谷 素顕、竹内 哲也、上山 智、赤﨑 勇、X 線その場観察 MOVPE による AlGaN/GaN ヘテロ構造評価、第 76 回応用物理学会秋季学術講演会シンポジウム、名古屋、2015323.
清木 良麻、小森 大資、池山 和希、伊奈 稔哲、小沼 猛儀、宮嶋 孝夫、竹内
哲也、上山 智、岩谷 素顕、赤崎 勇、X 線吸収微細構造測定による Al0.82In0.18Nの局所構造解析、第 76 回応用物理学会秋季学術講演会シンポジウム、名古屋、2015324.
林 貴文、永田 訓章、千賀 崇史、金山 亮介、岩山 章、岩谷 素顕、竹内 哲
也、上山 智、赤﨑 勇、電子線励起法による GaInN/GaN-MQW レーザ、第 76 回
応用物理学会秋季学術講演会シンポジウム、名古屋、2015
325.
宮嶋 孝夫、小森 大資、清木 良麻、伊奈 稔哲、新田 清文、鈴木 健太、竹内
哲也、宇留賀 朋哉、上山 智、岩谷 素顕、赤﨑 勇、XAFS 法を用いた GaN1-xSbx
混晶半導体中の Sb 占有位置評価、第 76 回応用物理学会秋季学術講演会シンポジ
ウム、名古屋、2015＜研究成果の公開状況＞（上記以外）
シンポジウム・学会等の実施状況、インターネットでの公開状況等
 第 6 回窒化物半導体結晶成長講演会（日本結晶成長学会ナノエピ分科会と共同主催）
開催日：平成 26 年 7 月 25 日～26 日 実施場所 名城大学 天白キャンパス
講演者：基調・特別・チュートリアル・招待講演 8 名、一般講演 63 名
掲載 HP：http://www.jacg.jp/jacg/japanese/nanoepi/2014.7.25-26/index.html
 2013 Meijo International Symposium on Nitride Semiconductors（主催）開催日：平成 25 年 11 月 29 日 実施場所 名城大学 天白キャンパス内
HP 掲載：http://nitride.meijo-u.ac.jp/iwaya/MSN/index-2013.html
講演者： 記念講演 2 名、招待講演 2 名（海外）
備考：本シンポジウムに長年貢献頂いた Bo Monemar 教授および Bernard Gil 教授に名城大学名誉博士号を授与した。
（参照：http://www.meijo-u.ac.jp/research/news/detail.html?id=8J94Jm）
 第 13 回赤﨑記念研究センターシンポジウム（協賛）
開催日：平成 25 年 11 月 29 日 実施場所 名古屋大学ベンチャー・ビジネス・ラボラトリー ３
階 ベンチャーホール
講演者：一般講演 3 名、招待講演 6 名
 2012 Meijo International Symposium on Nitride Semiconductors（主催）開催日：平成 25 年 2 月 28 日、実施場所 名城大学 天白キャンパス内
講演者： 招待講演 5 件
HP 掲載：http://nitride.meijo-u.ac.jp/iwaya/MSN/index-2012.html
 第 12 回赤﨑記念研究センターシンポジウム（協賛）
開催日：平成 25 年 2 月 27 日 実施場所 名古屋大学東山キャンパス野依記念学術交流館

（様式 2）
法人番号プロジェクト番号

231029S1201019

講演者：一般講演 6 件、招待講演 6 件
 2015 Meijo International Symposium on Nitride Semiconductors
開催日：平成 27 年 11 月 21 日（土）、実施場所 名城大学天白キャンパス内
講演者：一般講演１件、招待講演 3 件
 第 15 回赤﨑記念研究センターシンポジウム（協賛）
開催日：平成 27 年 11 月 20 日（金）、実施場所 名古屋大学
講演者：一般講演 8 件、招待講演 3 件
 The 4th International Workshop on LEDs and Solar Applications
開催日：平成 28 年 3 月 30 日（水）-31 日（木）、実施場所 名城大学天白キャンパス内
講演者：一般講演 3 件、招待講演 7 件

１４ その他の研究成果等
本研究拠点の整備にともない、多くの学生が大学院・修士課程および博士課程に進学し、博士課程修了学生 4 名（飯田・森下・安田・松井）および 80 名以上の大学院修士課程修了者を輩出した。これらの研究者は、国内はもとより海外（4 月よりサウジアラビアの KAIST で研究活動を実施予定）で活躍しており、この分野の発展に大きく貢献している。

（様式 2）
法人番号プロジェクト番号

231029S1201019

１５ 「選定時」及び「中間評価時」に付された留意事項及び対応＜「選定時」に付された留意事項＞
基盤形成の観点から研究成果に留意してほしい＜「選定時」に付された留意事項への対応＞
本プロジェクトでは、名城大学の窒化物半導体に関する研究者を中心に、結晶成長・デバ
イスプロセスの革新が可能な体制を構築し、それによって革新的なデバイスを実現できる基盤を形成することを目的に研究を進めてきた。特に対外発表を重視し、多数の学会や論文発表が形成できることに留意して研究を進めてきた。
また、グローバル化に向けて多くの海外研究者と若手研究者の交流を進めてきており、4名の博士課程・修士課程学生の輩出など拠点を形成する上で重要な点に留意して研究を進めてきた。さらに本プロジェクト終了後にも多くの国家プロジェクトを継続的に実施する予定であり、研究基盤形成という観点から十分に多くの成果を残したと考えられる。

＜「中間評価時」に付された留意事項＞
特になし＜「中間評価時」に付された留意事項への対応＞
特になし

Properties of the main Mg-related acceptors in GaN from optical and structural studies
B. Monemar, P. P. Paskov, G. Pozina, C. Hemmingsson, J. P. Bergman, S. Khromov, V. N. Izyumskaya, V.
Avrutin, X. Li, H. Morkoç, H. Amano, M. Iwaya, and I. Akasaki

Citation: Journal of Applied Physics 115, 053507 (2014); doi: 10.1063/1.4862928
View online: http://dx.doi.org/10.1063/1.4862928
View Table of Contents: http://aip.scitation.org/toc/jap/115/5
Published by the American Institute of Physics

JOURNAL OF APPLIED PHYSICS 115, 053507 (2014)

Properties of the main Mg-related acceptors in GaN from optical andstructural studies
B. Monemar,1 P. P. Paskov,1 G. Pozina,1 C. Hemmingsson,1 J. P. Bergman,1 S. Khromov,1
V. N. Izyumskaya,2 V. Avrutin,2 X. Li,2 H. Morkoç,2 H. Amano,3 M. Iwaya,4 and I. Akasaki4
1

Department of Physics, Chemistry and Biology, Link€
oping University, S-581 83 Link€
oping, Sweden
Department of Electrical and Computer Engineering and Physics Department, Virginia Commonwealth
University, Richmond, Virginia 23284, USA
3
Department of Electrical Engineering and Computer Science, Nagoya University, Chikusa-ku, Nagoya,464-8603, Japan
4
Department of Materials Science and Engineering, Meijo University, 1-501 Shiogamaguchi, Tempaku-ku,
Nagoya 468-8502, Japan
2(Received 25 October 2013; accepted 9 January 2014; published online 3 February 2014)
The luminescent properties of Mg-doped GaN have recently received particular attention, e.g., in thelight of new theoretical calculations, where the deep 2.9 eV luminescence band was suggested to be
the main optical signature of the substitutional MgGa acceptor, thus, having a rather large binding
energy and a strong phonon coupling in optical transitions. We present new experimental data onhomoepitaxial Mg-doped layers, which together with the previous collection of data give an
improved experimental picture of the various luminescence features in Mg-doped GaN. In n-type
GaN with moderate Mg doping (<1018 cm3), the 3.466 eV ABE1 acceptor bound exciton and theassociated 3.27 eV donor-acceptor pair (DAP) band are the only strong photoluminescence (PL)signals at 2 K, and are identified as related to the substitutional Mg acceptor with a binding energy of
0.225 6 0.005 eV, and with a moderate phonon coupling strength. Interaction between basal planestacking faults (BSFs) and Mg acceptors is suggested to give rise to a second deeper Mg acceptor
species, with optical signatures ABE2 at 3.455 eV and a corresponding weak and broad DAP peak at
about 3.15 eV. The 2.9 eV PL band has been ascribed to many different processes in the literature. It
might be correlated with another deep level having a low concentration, only prominent at high Mg
doping in material grown by the Metal Organic Chemical Vapor Deposition technique. The origin of
the low temperature metastability of the Mg-related luminescence observed by many authors is here
reinterpreted and explained as related to a separate non-radiative metastable deep level defect, i.e.,
C 2014 AIP Publishing LLC. [http://dx.doi.org/10.1063/1.4862928]
not the MgGa acceptor. V

I. INTRODUCTION

Mg is so far the only acceptor dopant that can be used
for p-doping in III-nitride based devices, such as light emitting diodes (LEDs) and laser diodes (LDs), for visible andultraviolet (UV) wavelengths. Although much work has been
done in an attempt to clarify the electronic and optical properties of Mg acceptors, there still remain open questions
regarding the interpretation of experimental data on the properties of Mg-doped III-nitride materials. This holds true evenfor GaN, the most developed III-nitride compound in termsof crystal perfection and doping control. In this paper, weshall concentrate on the optical properties of Mg-doped
GaN, specifically on luminescence with laser excitation(photoluminescence (PL)) and with electron excitation (cathodoluminescence (CL)). Such optical data give the mostdetailed information on the energy level structure of acceptors via their characteristic spectra, and also on the related
recombination mechanisms of carriers and excitons.
The main PL (or CL) spectra related to acceptors insemiconductors are the acceptor bound exciton (ABE) anddonor-acceptor pair (DAP) spectra, as reviewed for GaN in
Ref. 1. These PL signatures for Mg-related acceptors were0021-8979/2014/115(5)/053507/8/$30.00

established only in recent years,2–4 with the ABE1 peak at
3.466 eV in unstrained GaN at 2 K and the corresponding
DAP peaking at about 3.27 eV. The latter PL emission wasstudied in detail already four decades ago5,6 and established
as a DAP emission, but the connection with the Mg acceptor
was not definitely known at the time. An unusual feature
with Mg acceptors is that there are two ABEs present at highMg doping, the additional one being a broader ABE2 peak at
about 3.455 eV at 2 K.4 Another puzzling observation is theinstability of the Mg related spectra vs. UV excitation orelectron excitation in the low temperature region for p-typematerials, manifesting itself as the emission intensities going
down with excitation time.4,7–9 This effect is particularly
strong before the Mg-activation anneal of the samples, andhas raised questions whether the above-mentioned opticalsignatures are related to the substitutional MgGa acceptor or
possibly to an unstable Mg-related complex.4
At lower photon energies additional spectra in
Mg-doped GaN are noted. A broad peak at about 2.9 eV iscommonly observed at high Mg doping concentrations in
Metal-Organic Chemical Vapor Deposition (MOCVD)
grown samples.10,11 This has most often been assigned tothe presence of Mg–nitrogen vacancy (Mg-VN) deep donor

115, 053507-1

C 2014 AIP Publishing LLC
V

053507-2

Monemar et al.

J. Appl. Phys. 115, 053507 (2014)

complexes in the material.10 In the deep red spectral region
there is a broad band observed at about 1.8 eV, suggested as
related to a deep Mg-related complex defect.12–14
In this paper, we will present additional PL and CL dataobtained from MOCVD-grown GaN:Mg samples on low
defect density bulk GaN substrates. These data allow a moredefinite identification of the above-mentioned characteristic
ABE and DAP spectra for GaN:Mg. Furthermore, the reason
for the apparent presence of two acceptor species will be discussed in terms of interaction between Mg acceptors andstructural defects induced by the acceptor doping. For comparison, we also include new data from Mg-doped GaN samples grown by Halide Vapor Phase Epitaxy (HVPE).
II. SAMPLES AND EXPERIMENTAL PROCEDURE

The MOCVD samples, used in this work, were grownon thick (>300 lm) bulk GaN templates. Some of these thick
GaN substrates were grown with HVPE at Link€oping
University, and removed from the sapphire substrates with alaser lift-off technique, as reported previously.15 Other bulk
c-plane GaN substrates were supplied by Furukawa Co.,
Ltd., Japan. The c-plane Mg doped MOCVD samples weregrown at Meijo University as well as at Bremen University,
details of the growth procedure can be found in previous
papers.4 The m-plane GaN:Mg samples were grown at
Virginia Commonwealth University, on m-plane bulk GaN
substrates provided by Kyma, Inc.16 Some HVPE grownthick (about 11 lm) GaN:Mg layers on sapphire were provided by Oxford Instruments, the growth conditions aredescribed in Ref. 17.
Thermal annealing at Meijo University and Link€oping
University was done at 800  C for 10 min in N2 atmosphere.
At Bremen University Rapid Thermal Annealing (RTA) was
done during 1 min at 800  C. Stationary PL spectra were
measured with above band gap continuous wave UV excitation (laser photon energy of 4.65 eV) at temperatures from2 K to 300 K, and detected by a UV enhanced liquid nitrogencooled Charge Couple Device (CCD). PL transient measurements were done using femtosecond pulses from an amplified and frequency tripled Ti:sapphire laser (frequency
250 kHz and photon energy 4.65 eV), and detected with a
UV sensitive Hamamatsu streak camera with a slow sweep
unit. Cross-sectional Transmission Electron Microscope(TEM) analysis was done with a high resolution FEI Tecnai
G2 200 keV Field Emission Gun (FEG) instrument. CL
spectra were measured using a MonoCL4 system integrated
with a LEO 1550 Gemini scanning electron microscope(SEM) and equipped with a liquid-He-cooled stage forlow-temperature experiments. A fast CCD detection systemor a Peltier cooled photo-multiplier tube (PMT) was used forspectral acquisition.
III. EXPERIMENTAL RESULTS
A. Photoluminescence spectra in the near bandgap
regionThe main features of Mg-related bound exciton (BE)spectra in GaN were discussed in Ref. 4. In Fig. 1(a), we

FIG. 1. Low-temperature near-bandgap spectra of nominally undoped (a)and Mg-doped ([Mg] ¼ 3  1018 cm3) (b) homoepitaxial layers grown by
MOCVD on m-plane bulk GaN substrates.

show new PL data that illustrate the difference in PL spectra
at 2 K for nominally undoped GaN and highly Mg-doped
GaN, in this case for m-plane samples. Separate Secondary
Ion Mass Spectroscopy (SIMS) data on other samples indicatethat for a nominally undoped GaN layer grown on a bulk GaN
substrate the residual Mg concentration can be at most about
1  1017 cm3 at the interface between the layer and the substrate, and decreases to low 1016 cm3 upon further growth.
This Mg contamination can be ascribed to the well-known Mg
memory effect in MOCVD growth,18 and varies from sample
to sample. The PL of the nominally undoped sample in Fig.
1(a) does indeed show the two silicon and oxygen related donor BEs (DBEs) at 3.472 eV and 3.471 eV, respectively, butalso clearly the Mg-related ABE1 peak at 3.466 eV (in fact,
consisting of two separate peaks, as may be expected for
ABEs in GaN (Ref. 19)). We note that in this spectrum there
is no trace of the second ABE2 peak previously reported formore highly doped GaN:Mg.4 Likewise, there is no trace ofother sharp line spectra related to Mg doping, previously
reported in the region 3.35 eV to 3.42 eV.20 In n-type samples
with low Mg-doping, the metastable properties of the
Mg-related luminescence are also absent. Another sample
with a controlled Mg doping of [Mg] ¼ 1  1017 cm3 shows
very similar spectra as in Fig. 1(a).
In the m-plane sample doped with Mg to a concentrationof 3  1018 cm3 there is a broad ABE2 peak (Fig. 1(b)), inagreement with earlier data on c-plane samples.4 This ABE2

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FIG. 3. Near-bandgap PL spectra of as-grown HVPE GaN:Mg layer ([Mg]2  1019 cm3) at different temperatures.
FIG. 2. PL spectra of c-plane MOCVD homoepitaxial GaN:Mg([Mg] ¼ 1  1019 cm3) measured at different temperatures after annealing.

peak has a very similar polarization behavior as the ABE1
peak (dominantly E?c), indicating that both are somehow
related to a substitutional MgGa acceptor. Moreover, there isalso a series of rather strong peaks in the region
3.35–3.42 eV, in addition to the broadened phonon replicas
related to free and bound excitons present in the same energy
range. The origin of these lines is different basal plane stacking faults (BSFs), as discussed in Ref. 21.
Shown in Fig. 2 is the temperature dependence of PL fora c-plane MOCVD sample, with a Mg concentration of
approximately 1  1019 cm3. The sample is annealed, whichin this case enhances the relative intensity of the ABE2
peak.4 The spectral PL development with temperature isvery much the same for the two peaks ABE1 and ABE2, andthey both disappear at about 50 K, leaving behind only theDBE peak in the near bandgap region. Note that the narrower
ABE1 peak is observed at the high energy flank of the broad
ABE2 peak. These temperature dependent data are a supportof the ABE nature of the broad ABE2 peak, which due to thespectral broadening cannot be established with standard
spectroscopic means (e.g., via resolving the electronic substructure of the no-phonon line19). We note that the DBE PL
features in GaN are stable up to much higher temperatures,e.g., 150 K in samples with low doping.22
In order to obtain a complete experimental picture, thespectra from an 11-lm-thick HVPE grown GaN:Mg layer onsapphire was studied as well, as shown in Fig. 3. The Mg
concentration was not exactly known in this sample, but thehole concentration was measured as 2  1018 cm3 at room
temperature, which is consistent with an Mg concentrationof about 2  1019 cm3.17 The spectra clearly show the two
Mg related ABEs with a similar appearance as in the
MOCVD Mg-doped sample in Fig. 1(b). Previous work inliterature also shows the presence of both the ABE1 and theABE2 peak in Molecular Beam Epitaxy (MBE) grown
GaN:Mg at high doping concentrations, although the identityof the peaks was incorrectly labeled.23 Therefore, the presence of two Mg-related acceptor bound exciton near bandgap
PL signatures in GaN appears to be universal.
B. The DAP spectrum at 3.27 eV

A common feature in all acceptor related luminescence
spectra in semiconductors is that the ABE spectra in the near
bandgap region are accompanied by DAP spectra at lower
energies, i.e., there is one DAP emission and one ABE peak foreach particular acceptor species. For low Mg doping levels, therelation between the ABE1 peak at 3.466 eV and the 3.27 eV
DAP peak (with accompanying LO phonon replicas at lower
energies) has been established in the literature.2,3 For Mg doping at concentrations [Mg] < 1018 cm3 these are the only
acceptor related PL features in this energy range, and therefore
have to be assigned to the substitutional Mg acceptor.
For higher Mg concentrations, the connection between the
ABE2 peak and a weak broad background under the low
energy part of the 3.27 eV peak has also been suggested.4 In
order to shed more light on this, we have carried out new time
resolved PL measurements in the DAP region for an Mgdoped sample with [Mg] ¼ 2  1019 cm3, see Fig. 4. Fromthese data, it is obvious that at longer delay times (several ls) a
broad spectrum peaking at about 3.15 eV is dominant, which isthen suggested to correspond to a broad DAP emission related
to the A2 acceptor with a larger binding energy, consistent
with the larger ABE2 binding energy (compared to ABE1)
observed in the near bandgap spectra. A longer decay time fordeeper DAP emissions is a general observation in semiconductors.1 A discussion on the suggested interpretation of the corresponding deeper acceptors in terms of MgGa acceptors
interacting with structural defects is presented below.
C. Deeper Mg-related spectra

At higher Mg doping levels, in the range of 1020 cm3 for
MOCVD samples, a strong broad PL band at about 2.9 eV is

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FIG. 4. Time-resolved PL spectra of m-plane MOCVD homoepitaxial
GaN:Mg ([Mg] ¼ 2  1019 cm3) obtained with a time delay of 1 ls between
each curve. (The repetition rate of the femtosecond laser excitation was
250 kHz.)

noted and sometimes dominates the PL spectra.10–14 We showsuch spectra in Fig. 5 for MOCVD grown samples, measured
under different excitation conditions. In the sample with moderately high doping (1.5  1019 cm3) clearly both the 3.27 eV andthe 2.9 eV emission band are seen separately, while in the highly
doped sample (1  1020 cm3) only the 2.9 eV emission is present. This shows that these are two separate emissions, and thatthey are only moderately affected by potential fluctuations in thematerial.12,24 The connection to Mg-VN complexes often suggested for the 2.9 eV emission (see, e.g., Ref. 10) also seems lessprobable, since such complexes are expected to have a low concentration after annealing.25–27 The deep 2.9 eV emission seems
to be strong only for highly doped MOCVD grown GaN:Mg,
and not observable (or at least very weak) for the corresponding
doping in MBE (Refs. 23 and 28) or HVPE grown materials(see, Fig. 6). This situation presumably reflects the very different
growth conditions in these three growth techniques.

FIG. 5. Low-temperature spectra of c-plane MOCVD homoepitaxial
GaN:Mg layers measured at different excitation powers: (a) [Mg]
1.5  1019 cm3 and (b) [Mg] 1  1020 cm3.

IV. DISCUSSION
A. Tentative model for the second deeper Mg related
acceptor responsible for ABE2

Acceptor bound excitons at high doping level have previously been studied for other compound semiconductors
like InP and ZnTe.29,30 In these cases, a substantial broadening of the ABE peak (corresponding to ABE1 in our case) isobserved; it broadens out towards lower energy and there isalso some downshift in the mean peak position. The broadening on the low energy side is explained as an effect of thedispersion of the distances in the lattice for interacting
acceptor pairs binding the excitons at high doping.29,30 Our
observations for the Mg acceptors in GaN with two distinct

FIG. 6. Low-temperature bound-exciton and DAP spectra of as-grown andannealed HVPE GaN:Mg layer ([Mg] 2  1019 cm3). (The modulation of
the spectra is due to the interference effect.)053507-5

Monemar et al.

ABE peaks (ABE1 and ABE2) over the entire doping range,i.e., say 1  1018 cm3 to 2  1019 cm3, are clearly different
from the above traditional description of ABEs at interacting
acceptors in the III-V and II-VI compounds. There is noappreciable shift of the main ABE1 acceptor peak withincreased Mg doping, and the second ABE2 peak is well separated from ABE1 over this range of concentrations. So weconclude that the two ABE peaks correspond to two different
acceptor species caused by the Mg doping. At an Mg concentration above about 2  1019 cm3, the ABEs interact
strongly so that a Mott-like transition occurs and excitons(ABEs) are no longer bound to the acceptors, thus the near
bandgap ABE spectra disappear.
The unusual situation that two different acceptor levels
are created by one acceptor dopant in the case of Mg-doped
GaN has prompted new structural investigations of
GaN:Mg.31,32 The observation that the second acceptor responsible for the ABE2 signature is only observed in highly
doped GaN:Mg suggests the possibility of a correlation withstructural defects, that have also been observed in the
GaN:Mg samples studied here.32 These structural defects
appear only significantly for [Mg] > 1018 cm3, and are dominantly BSFs in the case of homoepitaxial material.31,32
Some of these SF defects are very small in the c-plane samples (a couple of nm in size), and those could also be pyramid features, as reported for GaN:Mg grown on sapphire.33
These defects are presumably induced during growth by the
introduction of Mg atoms in the lattice.32
The observed properties of the A2 acceptor discussed
here may be understood in a similar way as for the BSFrelated PL in n-type GaN.34 In that case, it was argued thatthe localization of electrons (and subsequently excitons) to
BSFs in n-GaN is mainly induced by donor potentials in thevicinity of the BSF.34 In the case of acceptor doping theacceptor potential is much more localized, within <1 nm,
meaning that the acceptor hole may not easily delocalize tothe BSF plane, but will distort the local potential at the BSF,
thus assisting the BSF hole (and exciton) localization, so thatthe BSF luminescence is observed.32 The acceptor will, in
turn, be perturbed by the nearby BSF potential mainly governed by the spontaneous polarization field,21 positioned on
the average a few nm away in p-GaN with an Mg doping of1019 cm3. This perturbed potential should be sufficiently
different from an unperturbed Mg acceptor to give rise to theA2 acceptor and its characteristic ABE2 spectrum.
We note that there is independent evidence for the stimulation of BSF luminescence by acceptor doping, as recently
observed in a series of n-type m-plane homoepitaxial GaN
samples.35 Also in our m-plane Mg-doped samples there arestrong BSF-related spectra in the range 3.3–3.42 eV, apparently stimulated by the Mg doping (Fig. 7). These BSF PL
spectra are different from the ABE2 peak, again emphasizingthat the A2 acceptor is likely to be associated to a specific siteof the Mg acceptor in relation to the BSF plane, i.e., not randomly distributed acceptors. The microscopic geometry of theA2 acceptor center remains to be revealed by future studies.
The PL spectrum of the unperturbed A1 acceptor has theappearance expected for an isolated Mg acceptor.19 The complex defect aggregate (Mg acceptor plus BSF), on the other

J. Appl. Phys. 115, 053507 (2014)

FIG. 7. Low temperature CL spectra of m-plane MOCVD homoepitaxial
GaN:Mg ([Mg] ¼ 1  1019 cm3).

hand, is suggested to explain the additional binding energy(ABE2 is deeper than ABE1), and the broad lineshape ofABE2 (see, Figs. 1(b), 2, and 3). The peculiar lineshapeobserved for ABE2 may originate from a strong transfer process for excitons between different sites (from shallower BEs
to deeper ones). Such excitonic transfer has been observed previously in other III-V and II-VI semiconductors and explainedin terms of acoustic phonon assisted exciton tunneling.36,37
Transient PL spectroscopy on a highly doped GaN:Mg sample
clearly demonstrates decay processes dominated by such transfer, see Fig. 8. The range of acoustic phonons in GaN has been
determined to cover phonon energies up to above 30 meV(Ref. 38) therefore the ABE2 peak at 3.455 eV is well within
the range of a transfer process involving the ABE1 at 3.466 eV
as the main original state. The intensity profile of ABE2 has ahigh energy portion, which is, continuously rising from ABE1
all the way down to the ABE2 peak. The position of ABE2 israther well defined in all samples, and the acceptor binding
energy (as well as the ABE2 binding energy) should be determined by the specific configuration of the defect center, i.e.,the Mg acceptor–BSF combination.
The PL decay traces in Fig. 8(a) clearly show the fast
initial decay of the shallower DBE and ABE1 signal, much
faster than the radiative decay times observed in samples
with lower doping.39 In Fig. 8(b), a time resolved development of the ABE1–ABE2 pair of peaks is shown. These data
show directly the spectral shift of the PL signal from theABE1 position to ABE2, mostly during the first 100 ps. Thisis naturally explained as exciton transfer, as discussed above.
The broad ABE2 PL peak has a very specific dependence on excitation intensity, i.e., at low excitation levels thepeak is just barely seen, while it increases superlinearly athigh excitation levels (Fig. 9). This is again consistent withthe exciton transfer processes. At low excitation levels there
are few neutral acceptors and ABEs, and therefore the distance of an ABE site to another neutral acceptor is long. Thismeans a low probability for transfer because the overlap matrix element is very small. With increasing excitation the053507-6

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FIG. 9. Excitation intensity dependent PL spectra of m-plane MOCVD
homoepitaxial GaN:Mg ([Mg] ¼ 1  1018 cm3).

FIG. 8. (a) PL decay curves taken at different photon energies for c-plane
MOCVD homoepitaxial GaN:Mg ([Mg] ¼ 1.5  1019 cm3) and (b)time-resolved PL spectra of the same sample.

occupation probability of both neutral acceptors and neutral
ABEs (ABE1) increases, and as a result the distance between
ABE1 sites and neutral A2 acceptors becomes favorable fortransfer.
B. The 3.27 eV DAP emission and blue PL band at
2.9 eV

While the relation of the 3.27 eV PL spectrum to the
MgGa acceptor seems to be well established experimentally,
the identity of the blue PL band peaking at about 2.9 eV is
confusing within the confines of the available data and theirinterpretation in the literature. As stated above, the early
investigations focused on a DAP recombination model witha deep VN-Mg donor complex together with a substitutional
Mg acceptor.10 This DAP model is only consistent with ahigh concentration of these donor complexes (>1019 cm3),
while the investigations that followed estimated the actual
concentration of these donors to be as low as 1017 cm3
before annealing and much lower after anneal.25,26 Further,
the Optically Detected Magnetic Resonance (ODMR) investigations of GaN:Mg failed to show any evidence for deep
donors.40 Recently, another interpretation of this blue PL asa DAP involving shallow donors and the substitutional MgGa
acceptor was presented, based on theoretical calculations
indicating that the Mg acceptor should be deeper than previously believed and have a strong phonon coupling.41 This
interpretation also means that the 3.27 eV DAP is not related
to the MgGa acceptor, but to an Mg-H complex. The blue PL
is only strong in MOCVD grown GaN at high Mg doping,
however, while the same material grown with MBE or
HVPE does not show the blue PL peak, in spite of the factthat a similar number of MgGa acceptors should be present.
There are many early papers demonstrating the DAP natureof the 3.27 eV PL.5,6 A most relevant recent paper on low
Mg doped MBE samples demonstrate clearly both the DAP
and the free-to-bound (FB) transition, at different temperatures, confirming the binding energy of the Mg acceptor as
225 6 5 meV.42 An argument against the interpretation of
the 3.27 eV emission as being related to a neutral Mg-H
complex is that this complex (from theoretical results41) is
only expected to be found in p-type samples with a high Mg
doping, while in reality the 3.27 eV PL is also prominent in
n-type GaN. It is, in fact, the only Mg-related PL emission inn-type GaN, and dominates the PL spectrum in all Mg-doped
samples, except in very highly doped MOCVD samples
where the 2.9 eV band is strong. This behavior is expected if
the 3.27 eV DAP involves the MgGa acceptor.
Returning to the deeper 2.9 eV band there is some evidence from ODMR data for a DAP recombination, since ashallow donor signal is observed for detection in thispeak.43,44 Thermal quenching of the emission is observed
near room temperature with an activation energy of
0.3–0.4 eV.45 This would be consistent with an acceptor
level of similar binding energy, deeper than for the MgGa
acceptor. The oscillator strength for the deeper 2.9 eV transition is naturally weaker than that for the 3.27 eV PL line,
explaining the frequently observed spectral upshift in thisspectral range upon increasing excitation density.
The annealing studies in H plasma reported in Ref. 11 areinteresting. It is shown that the 2.9 eV PL seems to be Hrelated, and the corresponding defect does not noticeably
affect the hole concentration. The authors explain their data
regarding a DAP process with an H-related deep donor (binding energy 0.37 eV) and the shallow Mg acceptor, similar toother work.45 An alternative model would be a H-related

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complex that gives rise to the 2.9 eV PL line, but has a concentration much lower than the [Mg] atomic concentration,and thus does not affect the hole concentration in p-GaN.
C. Instabilities of acceptor related luminescence
spectraAn important property that needs an explanation is theinstability of the PL or CL spectra (ABE as well as 3.27 eV
DAP spectra) during prolonged excitation with above
bandgap UV light or energetic electrons.4,7–9 This is particularly important for MOCVD material which usually contains
an appreciable concentration of H. In our previous work, it
was suggested that this property could be interpreted as evidence for instability of the acceptor responsible for the
ABE1 and the 3.27 eV spectral features.4 Since our later optical investigations discussed in this paper are only consistent
with a model where the ABE1 and the 3.27 eV PL features
are related to the substitutional Mg acceptor, we have tested
alternative models to explain these instabilities. In Fig. 10,
we show time sequential PL spectra of an m-plane Mg doped
sample, obtained at time intervals of 10 s during continuous
UV excitation. It is obvious that in this case, the dominantfeature of this short term instability is a rapid degradation of
the PL efficiency in the entire near bandgap region, evidence
for a rapid nonradiative process induced by the UV excitation. Such a process is most naturally explained by the presence of a defect possessing instable properties, acting as a
non-radiative shunt path for the near bandgap luminescence.
The involvement of another defect (i.e., not primarily
the Mg acceptor) in the instability process is further indicated by time-resolved PL data of the DAP emission at
3.27 eV (see, Fig. 11). This figure shows that upon long term
excitation there is a clear downshift in energy of the 3.27 eV
DAP peak, evidence for a decrease of the hole occupation of
the Mg acceptor.46,47 This decrease means that the quasi
Fermi level is raised in the bandgap during the long term excitation in this experiment, which is understood if anotherdeeper (nonradiative) defect level is active in the recombination process, capturing photoexcited holes to a metastablestate, which decreases the hole occupation in the Mg
acceptor. This nonradiative defect is then directly

FIG. 10. Low-temperature spectra of m-plane MOCVD homoepitaxial
GaN:Mg ([Mg] ¼ 3  1018 cm3) obtained at time intervals of 10 s duringcontinuous UV excitation.

FIG. 11. Low-temperature spectra of c-plane MOCVD homoepitaxial
GaN:Mg ([Mg] ¼ 1.5  1019 cm3) measured upon long term continuous
UV excitation showing a clear red shift of the DAP emission.

responsible for the long term instability properties. To reveal
specific properties of the defect responsible for the long term
instability, more experiments to obtain Deep Level Transient
Spectroscopy (DLTS) data are needed.
As mentioned above, this idea with an independent deeplevel defect responsible for the instability allows the interpretation of the strongest features in PL and CL data of
Mg-doped GaN, the ABE1 peak and the 3.27 eV DAP, asconnected to the regular substitutional Mg acceptor, a simpler model than the one discussed in Ref. 4. Unfortunately,electrical measurements to reveal properties of the metastable defect are very difficult in the low temperature range
involved here, and very limited such data for Mg-doped
p-GaN are found in the literature.48 Metastable defects arealso reported for n-GaN, suggested to be VGa-related.49,50
Such defects related to VGa would be unlikely in p-GaN(Ref. 51) therefore, more work is needed to explore the identity of the metastable defect in Mg-doped GaN.
V. CONCLUSIONSWe discuss new and previous optical data for the Mg
acceptor in GaN, in an attempt to arrive at the simplest possible model that can explain all the data in aggregate.
Intriguing properties discussed in previous work were the existence of two different acceptor signatures in low temperature PL spectra for bound excitons (ABE1 at 3.466 eV and
ABE2 at 3.455 eV), and the corresponding DAP1 at 3.27 eV
and the weaker DAP2 at about 3.15 eV. Since at low Mg
doping (n-type samples) only ABE1 and DAP1 are observed,
the above-mentioned spectra should be related to the substitutional Mg acceptors, i.e., no Mg-H related complexes aspreviously suggested in relation to the observed instabilities4
are needed to explain these optical data. The second deeper
set of broadened spectra ABE2 and DAP2 appear at higher
doping ([Mg] > 1018 cm3), simultaneously, as structuraldefects (basal plane stacking faults) are introduced.32 The

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acceptors involved in the ABE2 and DAP2 peaks are therefore assigned to substitutional Mg acceptors perturbed by the
induced structural defects (BSFs). This correlation was madepossible in this work via the use of homoepitaxy on bulk
GaN substrates, thus avoiding the more complex spectral
behavior in samples grown on sapphire.52 The unusual unstable behavior in the optical spectra at low temperatures forhigh Mg doping may then be explained by the presence ofanother nonradiative recombination path via a deeper defectlevel which is metastable in this low temperature range(<300 K), and affects the PL of the Mg-related spectra via
the quasi-Fermi level under optical excitation. Since thedeeper radiative PL emission at 2.9 eV only occurs at highdoping levels in MOCVD samples (and is not regularly
observed in similarly Mg-doped MBE and HVPE samples),
it is implausible that this emission is related to the substitutional Mg acceptor, as recently suggested in theoreticalwork.41,53 It may relate to a deeper acceptor of low concentration, introduced primarily in MOCVD growth.
ACKNOWLEDGMENTS

We are grateful to S. Figge and D. Hommel at Bremen
University for some of the homoepitaxial MOCVD samples
used in this study and to A. Usikov at Oxford Instruments,
Oxford, UK for a p-type Mg-doped HVPE sample. The
Link€
oping group also acknowledges support from the K. A.
Wallenberg Foundation for the financing of spectroscopic equipment and the Li-Li-IFM Linneaus project for partial support.
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Laser lift-off technique for freestanding GaN substrate using an In droplet formed bythermal decomposition of GaInN and its application to light-emitting diodes,
Daisuke Iida , Syunsuke Kawai, Nobuaki Ema, Takayoshi Tsuchiya, Motoaki Iwaya, Tetsuya Takeuchi, Satoshi
Kamiyama, and Isamu Akasaki

Citation: Appl. Phys. Lett. 105, 072101 (2014); doi: 10.1063/1.4893757
View online: http://dx.doi.org/10.1063/1.4893757
View Table of Contents: http://aip.scitation.org/toc/apl/105/7
Published by the American Institute of Physics

APPLIED PHYSICS LETTERS 105, 072101 (2014)

Laser lift-off technique for freestanding GaN substrate using an In droplet
formed by thermal decomposition of GaInN and its application
to light-emitting diodes
Daisuke Iida,1,a) Syunsuke Kawai,1 Nobuaki Ema,1 Takayoshi Tsuchiya,1 Motoaki Iwaya,1Tetsuya Takeuchi,1 Satoshi Kamiyama,1 and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

2(Received 2 August 2014; accepted 8 August 2014; published online 18 August 2014)
We developed a laser lift-off technique for a freestanding GaN substrate using an In droplet
formed by thermal decomposition of GaInN. A combination of an In droplet formed by thermal
decomposition of GaInN during growth and a pulsed second-harmonic neodymium-doped yttrium
aluminum garnet laser (k ¼ 532 nm) realized the lift-off GaN substrate. After laser lift-off of the
GaN substrate, it was used to achieve 380 nm ultraviolet light-emitting diodes with light output
enhanced 1.7-fold. In this way, the light extraction can be improved by removing the GaN
C 2014 AIP Publishing LLC. [http://dx.doi.org/10.1063/1.4893757]
substrate. V

Group-III-nitride-based light-emitting diode (LED)technology appears to be an excellent candidate for solidstate platforms. Currently, high-brightness GaInN-based
visible LEDs have been commercialized.1–4 In particular, theexternal quantum efficiency of blue LEDs has reachedapproximately 83%.5 However, the ultraviolet (UV) LED
still has a low external quantum efficiency because of a lowinternal quantum efficiency, low free hole carrier concentration in the p-layer, and high internal absorption by the metalelectrodes and epilayer.6–10 The UV LED is also a major
focus in the research efforts of group-III-nitride-based LEDs.
High-efficiency UV-A LEDs are usually grown on sapphire substrates fabricated by a laser lift-off (LLO) techniqueand polishing/etching of the GaN epilayer after lift-off fromthe sapphire substrate.11 It is well known that of UV-A
LEDs are strongly dependent on the dislocation density of
the underlying layer. With the adoption of a high-quality
low-dislocation-density GaN substrate, the UV-A LEDs are
expected to have improved internal quantum efficiency as aresult of reduced nonradiative recombination.12,13 It is significantly important to reduce the dislocation density of theunderlying layer to realize high-performance devices. The
light extraction efficiency can also be improved by using thesurface texturing technique.14 However, growing UV-A
LEDs on GaN substrates is avoided because the GaN
substrate has a degradation factor of light extraction owing
to internal absorption loss due to the presence of impurities(e.g., oxygen and carbon) and their complexes.15,16 For therealization of high-efficiency UV-A LEDs grown on GaN
substrates, it is required to remove the GaN substrate. We
look to expect a LLO, which is one technique of overcoming
GaN absorption loss. According to several reports, the realization of LLO LEDs grown on sapphire substrates has
almost been achieved.17–20 A typical LLO technique enables
us to separate the sapphire substrate from the interface with
GaN using a high-energy pulsed UV laser (e.g., KrF pulsed
excimer laser; k ¼ 248 nm), because they have a hugea)E-mail: dft0tfi16@meijo-u.ac.jp0003-6951/2014/105(7)/072101/4/$30.00

difference in absorption energy. Thus, the GaN layer nearby
the interface with sapphire can strongly absorbed the UV
laser light, which induces a Ga droplet formed by the decomposition of GaN. However, it then leads to the degradation of
the quality of the active layer owing to the absorption of thehigh-energy pulsed UV laser light. On the other hand, a GaN
epilayer on the GaN substrate has almost the same absorption energy gap. Therefore, it is impossible to remove the
GaN substrate from the UV-A LED structure by the typical
LLO technique.
In this study, we developed a LLO technique for freestanding GaN substrate using adopting an In droplet layer
formed by thermal decomposition of GaInN during growth.
The In droplet layer is expected to absorb light at visible andinfrared wavelengths by surface plasmon resonance.21 Our
LLO technique enabled the separation of the GaN substratefrom the epilayer through the use of a pulsed secondharmonic neodymium-doped yttrium aluminum garnet(Nd:YAG) laser (k ¼ 532 nm) as a result of the absorption ofvisible light in the In droplet layer. The LLO LED had
improved electron-luminescence (EL) emission intensity
owing to the reduced absorption loss of the GaN substrate.
Specifically, our LLO technique is expected to be useful forthe reuse of the GaN substrate and the utilization of somedevices to realize low cost and high performance.
All the samples were grown by metalorganic vapor
phase epitaxy in a face-down 2 in.  3 wafer horizontal-flow
reactor. We prepared a high-quality freestanding GaN substrate. The threading dislocation density of the GaN substrate
was below 106 cm2, as characterized by cathodoluminescence (CL) at room temperature. First, we investigated the
crystalline quality of a 2-lm-thick highly Si-doped n-Al0.03
Ga0.97N layer grown on a GaN substrate with/without
removing the layer of Ga0.85In0.15N/GaN superlattices (SLs).
The samples with and without Ga0.85In0.15N/GaN SLs are
assigned as samples (A) and (B) in Fig. 1, respectively. It is
well known that GaInN is a very sensitive material withthermal stability.22 After the growth of Ga0.85In0.15N/GaN
SLs, they were transformed into In droplets by the thermal

105, 072101-1

C 2014 AIP Publishing LLC
V

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Iida et al.

FIG. 1. Schematic view of n-Al0.03Ga0.97N layers (a) with and (b) without
Ga0.85In0.15N/GaN SLs grown on GaN substrate. Ga0.85In0.15N/GaN SLs
were transformed into In droplet layer by thermal decomposition of
Ga0.85In0.15N during n-GaN and n-Al0.03Ga0.97N growth.

decomposition of GaInN during n-GaN and n-Al0.03Ga0.97N
growth at 1040  C. The characterization of an as-grownhighly Si-doped n-Al0.03Ga0.97N layer with/without the In
droplet layer grown on a GaN substrate by X-ray diffraction(XRD) (Regaku: Smart Lab), transmittance (Shimadzu: UV2500PC), and CL (Hitachi: SU-70) revealed the crystal structure, optical property, and crystalline quality, respectively.
Second, we fabricated an UV-A LED with a GaInN multiple quantum well (MQW) active layer grown on a GaN
substrate. Figure 2 shows the fabrication steps of LLO andcontrol LEDs. The LLO LED consisted of 8 periods of
Ga0.85In0.15N/GaN SLs in order to remove the GaN substrate, a 1-lm-thick Si-doped n-GaN layer, 2-lm-thickhighly Si-doped n-Al0.03Ga0.97N layer, 10 periods of Ga0.97
In0.03N(2 nm)/GaN(2 nm) SLs, 5 periods of Ga0.95In0.05N(6 nm)/GaN(12.5 nm) MQWs, a 20-nm-thick Mg-doped
p-Al0.13Ga0.87N electron-blocking layer, and a 120-nm-thickMg-doped p-GaN contact layer. Actually, the Ga0.85In0.15N/
GaN SLs for the removing layer were almost completelytransformed into In droplets by the thermal decompositionof GaInN with increasing temperature to 1040  C during
n-GaN and n-Al0.03Ga0.97N growth. After exposure of theAppl. Phys. Lett. 105, 072101 (2014)

n-Al0.03Ga0.97N layer by inductively coupled plasma etching,a Ag(150 nm)/ITO(30 nm) contact was deposited on both nAl0.03Ga0.97N and p-GaN layers. Bare LED chips of
500  600 lm2 size were fabricated as flip-chip structures.
After mounting bare LED chips on submounts with Au
bumps, an epoxy resin was applied to fill the gap between
the LED chip and the submount.20 The LLO setup consistsof a pulsed second-harmonic Nd:YAG laser (k ¼ 532 nm).
The irradiation energy density was 125 mJ/cm2. Afterremoving the GaN substrate, we etched the back surface toimprove the light extraction efficiency, using hot potassiumhydroxide (KOH) solution after LLO.12,20 The control LED
was composed of the 200-lm-thick GaN substrate without
the Ga0.85In0.15N/GaN SLs for removing GaN substrate.
Finally, light output power- and voltage-injection current(V-I, L-I) characteristics of each LED were measured using aluminescence characterization system (Otsuka Electronics:
MCPD-9800).
The X-ray diffractions of samples (A) and (B) are shownin Fig. 3. We found that Ga0.85In0.15N/GaN SLs were almost
destroyed by thermal decomposition during the growth of nGaN and n-Al0.03Ga0.97N at 1040  C. At the same time, it
included the metal phase of In (i.e., In droplet) dissociatedfrom GaInN. We confirmed the existence of the In droplet
layer in sample (A) from the (101) diffraction. This resultshows that the Ga0.85In0.15N/GaN SL removing layer wascompletely transformed to In droplets during growth.
We characterized the transmittances of samples (A) and(B), as shown in Fig. 4. Sample (A) exhibits absorption in awide range from violet to infrared wavelengths. Regarding
the X-ray diffraction results, it was found that light absorption was due to the In droplet layer. Insets of Fig. 4 show
images of the as-grown samples. In particular, sample (A)had already turned black. It was found that the In droplet
layer can absorb visible light by surface plasmon resonance.
This causes selective absorption of visible or infrared pulse
laser light by the In droplet layer, because neither the GaN
substrate nor the LED structure can absorb visible and infrared light. In this way, we found that removing of the GaN
substrate technique was possible using a visible or infrared
pulse laser. Then, there is no worry about the degradation of

FIG. 2. Schematic view of fabrication
steps of (a) LLO and (b) control LEDs.

072101-3

Iida et al.

Appl. Phys. Lett. 105, 072101 (2014)

FIG. 3. XRD of samples (A) and (B) with and without removing layer,
respectively. The diffracted X-ray at the (101) plane indicated the metal
phase of In in sample (A) with the removing layer.

the quality of the MQW active layer owing to the absorptionof high-energy laser light.
We observed threading dislocation densities of samples(A) and (B) by CL. From the CL results, threading dislocation densities of samples (A) and (B) were estimated to be
5  107 and 3  106 cm2, respectively. We found that sample (A) had an increased threading dislocation density. We
consider that after the decomposition of Ga0.85In0.15N/GaN
SLs during growth, the threading dislocations were induced
via the generation of misfit dislocation in the interface
between n-GaN and the In droplet layer. The LLO LEDs predict that the internal quantum efficiency will be degraded
compared with that of the control LEDs. However, the dislocation density was sufficiently low compared with that ofconventional GaN grown on a sapphire substrate, the typicalvalue is 2–5  108 cm2. The quality of sample (A) was
sufficient for use in high-performance UV-A LEDs.
The EL spectra of LLO and control LEDs without andwith the GaN substrate, respectively, are shown in Fig. 5.
The LLO LED indicated strong EL emission compared withthat of the control LED. The LLO LED showed enhanced
light output power even though a decrease in internal

FIG. 4. Transmittance as a function of wavelength. Sample (A) absorbs lightin the entire visible range. Insets show images of as-grown samples of different colors. Scale bars are 1 cm.

FIG. 5. (a) EL spectra of each LED at 50 mA. (b) Inset EL spectra are on thelog scale. Insets shown are EL of LLO and control LEDs.

quantum efficiency due to the degradation of the crystalline
quality of the underlying n-AlGaN epilayer was predicted.
We also found that the EL peak emission of LLO LED wasshifted to a short wavelength from 383 to 380 nm. We consider that the EL intensity and wavelength were different
because they were cut off at a shorter wavelength by the
GaN substrate, which acted as an absorber. After LLO process, the yellow luminescence of the deep level recombination can be reduced.15,16 The L-I and V-I characteristics of
UV LEDs are shown in Fig. 6. The light output power of the
LLO LED can be enhanced 1.7-fold at 50 mA comparedwith that of the control LED. V-I characterizations were thesame because the highly Si-doped n-Al0.03Ga0.97N layer isan extremely low-resistance material.23 These results confirm that we have realized the LLO of the GaN substrate.
Our LLO technique can overcome the absorption loss due tothe GaN substrate. It is very promising for improving theexternal quantum efficiency of UV-A LEDs.
We achieved the removing of the GaN substrate fromthe UV-A LED structure by a LLO technique. Ga0.85In0.15N/
GaN SLs were transformed into an In droplet layer by the

FIG. 6. L-I and V-I characteristics of LLO and control LEDs.

072101-4

Iida et al.

thermal decomposition of GaInN during growth. It was able
to absorb light in the visible range, which enabled the use of
the visible pulse laser. After the lift-off process, the UV-A
LEDs show EL intensity enhanced by 1.7-fold. Therefore, it
is very effective to improve light extraction by reducing thelight absorption loss. Therefore, UV-A LEDs can also begrown on the GaN substrate by the LLO technique.
This work was supported by the Program for theStrategic Research Foundation at Private Universities,
2012–2016.
1

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10Internal quantum efficiency enhancement of GaInN/GaN quantum-well structuresusing Ag nanoparticles,,
Daisuke Iida, Ahmed Fadil , Yuntian Chen , Yiyu Ou, Oleksii Kopylov, Motoaki Iwaya, Tetsuya Takeuchi,
Satoshi Kamiyama, Isamu Akasaki, and Haiyan Ou

Citation: AIP Advances 5, 097169 (2015); doi: 10.1063/1.4931948
View online: http://dx.doi.org/10.1063/1.4931948
View Table of Contents: http://aip.scitation.org/toc/adv/5/9
Published by the American Institute of Physics

AIP ADVANCES 5, 097169 (2015)

Internal quantum efficiency enhancement of GaInN/GaN
quantum-well structures using Ag nanoparticles
Daisuke Iida,1,2,3 Ahmed Fadil,2,a Yuntian Chen,2,4,b Yiyu Ou,2Oleksii Kopylov,2 Motoaki Iwaya,3 Tetsuya Takeuchi,3 Satoshi Kamiyama,3Isamu Akasaki,3,5 and Haiyan Ou2
1

Department of Applied Physics, Tokyo University of Science,
Katsushika, 125-8585 Tokyo, Japan2
Department of Photonics Engineering, Technical University of Denmark,2800 Lyngby, Denmark
3
Faculty of Science and Technology, Meijo University, 1-501 Shiogamaguchi Tempaku,
468-8502 Nagoya, Japan
4
School of Optical and Electronic Information, Huazhong University of Scienceand Technology, 430074 Wuhan, China
5
Akasaki Research Center, Nagoya University, Furo-cho Chikusa, 464-8603 Nagoya, Japan(Received 22 May 2015; accepted 16 September 2015; published online 24 September 2015)
We report internal quantum efficiency enhancement of thin p-GaN green quantumwell structure using self-assembled Ag nanoparticles. Temperature dependent photoluminescence measurements are conducted to determine the internal quantum efficiency. The impact of excitation power density on the enhancement factor is investigated. We obtain an internal quantum efficiency enhancement by a factor of 2.3 at
756 W/cm2, and a factor of 8.1 at 1 W/cm2. A Purcell enhancement up to a factorof 26 is estimated by fitting the experimental results to a theoretical model for theefficiency enhancement factor. C 2015 Author(s). All article content, except whereotherwise noted, is licensed under a Creative Commons Attribution 3.0 Unported
License. [http://dx.doi.org/10.1063/1.4931948]

I. INTRODUCTION

An efficient high-brightness solid-state lighting solution is of central industrial and social relevance in terms of reducing energy consumption and environmental benefits. The GaInN-material
based light-emitting diodes (LEDs) technology appears to be an excellent candidate in solid-state
lighting platform. Improving the LEDs performance within the green gap is challenging, and isalso a major focus in the research efforts of GaInN-based LEDs. The LEDs grown on c-planesapphire substrate have a large piezoelectric field in the quantum well (QW) due to quantum
confined Stark effect (QCSE).1,2 The large piezoelectric field is one of the most critical factors
limiting the internal quantum efficiency (IQE) in the green gap. In order to improve GaInN-based
LEDs both enhancement of the IQE and development of an efficient light extraction techniquesto ‘drag’ the photons out of the LED crystal are required.3,4 It is well known that spontaneousemission, including the recombination rate and the emission profile, can be modified by the nearby
photonic environments.5,6 The key concept underlying plasmon enhanced light emission is the local
density of optical states (LDOS), which accounts for the available number of modes per frequency
at a certain position. Recent progress in nanophotonics and nanotechnology opens possibilities ofengineering LDOS at nanometer scale and at highly advanced levels to achieve directional emission or enhanced light-matter interaction.7–14 Such a nanophotonic approach has been applied tolight emitting devices.15–22,24,25 The rationale behind is that the transfer of energy from carriers in
GaInN QW into localized optical plasmon supported by the metallic nanostructures will create ana afad@fotonik.dtu.dk
b yuntian@hust.edu.cn

2158-3226/2015/5(9)/097169/9

5, 097169-1© Author(s) 2015

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AIP Advances 5, 097169 (2015)

additional light emission channel called plasmonic channel. With a thin p-GaN layer plasmonicchannel will dominate the spontaneous emission and significantly enhance the radiative recombination rate of light sources. As the rate of emitted photons is enhanced compared to the undesired
non-radiative channels, which are unaffected by the plasmonic structures, we expect to overcome
the challenge on IQE. Subsequently, the radiation captured by the plasmonic structure needs to be
scattered into desired output angles by designed grating effects from periodic plasmonic structures
or by exploiting random scattering from disordered plasmonic structures. Indeed, the IQE measurement of blue emission GaInN QWs indicates that significant plasmonic coupling can be achieved
at approximately 6.8-folds when the distance of the active layer to the metals (D) is very small,i.e. ∼10 nm.3 The blue LEDs having a small D of 30 nm is reported to have enhanced optical output
power by 38% using the nanoparticles embedded in p-GaN.17 For a larger D of approximately
70 nm, the green emission output power based on the plasmonic coupling of the metals with active
layer has been improved by approximately 90-220% with contribution from both IQE and light
extraction efficiency.16 It is also very interesting to note that the photoluminescence (PL) intensityof the green GaInN QW has been improved by a factor of 4.8 using periodic nanocylinders witha 5 nm spacing layer.20 It has been shown that for very large D ∼200 nm, a 26% improvement of
the optical output power due to enhanced light scattering can be achieved without benefits fromplasmon enhanced IQE.19 Despite the aforementioned evidences on improved PL or electroluminescence (EL) emission, to the best of our knowledge, there is no quantitative study of how originalemitter IQE affects the enhancement obtainable through LSP-coupling between metal and emitter.
In this paper, we experimentally demonstrated the IQE enhancement of plasmon-based LED
epi-structures by depositing metallic nanoparticles on top of the GaN crystal using a low-cost solution. The IQE is estimated from low temperature PL measurements. We perform excitation power
density dependent measurements, which show the IQE can be substantially enhanced with properly
selected nanometallic structure. The paper is organized as follows: In section II, we outline thefabrication procedures of our LED structures and the metallic nanoparticles, and demonstrate themeasurement setup. In section III, we present and discuss our experimental results of PL and IQEenhancement. Finally, section IV concludes the paper.

II. METHODS

LED structures were grown by metalorganic vapor phase epitaxy (MOVPE) in a horizontal
flow reactor. C-plane sapphire substrates were thermally cleaned in a H2 ambient at approximately
1050 oC. After that, approximately 20-nm-thick low temperature (LT)-GaN buffer layer was deposited at 535 oC.23 After annealing process a 2-µm-thick GaN layer was grown at 1050 oC on theannealed LT-GaN nucleation layer. The final LED structure consisted of a 2-µm-thick n-GaN:Si
layer, a 10 periods of GaInN:Si (3 nm) / GaN:Si (2 nm) superlattice layer, a 5 periods of GaN:Si(11.5 nm) / GaInN (2 nm) QWs active region covered with 5-nm-thick GaN capping layer, a20-nm-thick p-GaN:Mg, and a 5-nm-thick p+-GaN:Mg p-contact layer. The distance of last QW
to LED surface was 30 nm. This LED structure had an optimized thin p-GaN layer for couplingof localized surface plasmon (LSP) to QWs, which is expected to enhance the light emission of
QWs with LSP coupling. Self-assembled Ag NPs were fabricated by electron-beam evaporation andthermal annealing. Various-thickness Ag thin films were deposited on top of the LED surface. The
samples were annealed at 200 oC for 30 min in N2 atmosphere. Ag thin film was transformed into
random nanostructures during thermal annealing process.
The observations of Ag NPs on top of the LED surfaces were conducted by scanning electron
microscope (SEM). Excitation power density dependent PL measurement was carried out to characterize the emission of GaInN/GaN QWs grown on c-plane sapphire substrate. The PL measurement
setup consisted of a continuous wave GaInN-based diode laser and a spectrometer. The excitation
laser operated at 405 nm and the output power density could be tuned from 1 to 756 W/cm2.
PL measurement was carried out with excitation and collection via the backside of the polishedsapphire substrate. The extinction spectrum of Ag NPs was extracted from a transmittance spectrum
measurement. Transmittance was measured from the sapphire to the top surface p-GaN side, while

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FIG. 1. SEM images of Ag NPs on LED surface after thermal annealing process. Ag deposition thicknesses are (a) 4 nm
(sample A), (b) 9.5 nm (sample B) and (c) 17 nm (sample C), respectively.

reflectance was measured from the sapphire side. The light source used for the measurements was awhite-light Xenon lamp.

III. RESULTS AND DISCUSSIONS
A. Self-assembled Ag NPs

The Ag NPs were fabricated using a low-cost approach, i.e. by thermal annealing of Ag thinfilms on the surface of LED structure with designed thin p-GaN layers. Fig. 1 shows plan-view SEM
images of Ag NPs with different Ag deposition thickness. Samples A, B and C are Ag films witha thickness of 4, 9.5 and 17nm, respectively, which are transformed into random NPs after thermal
annealing on top of the LED p-GaN surface. As the Ag film thickness increases, the size of Ag NPs
becomes larger, while the density of Ag NPs decreases. Ag NPs had a tendency of becoming more
irregular-shaped and with an increased size distribution as the thickness was increased.
Fig. 2 shows the extinction spectra of the three samples. PL emission spectrum of GaInN/GaN
5QWs LED without Ag NPs constituting a reference sample is also shown in Fig. 2. Considering
the extinction spectra, each sample shows a characteristic peak. For sample A the extinction peakis located at 482 nm, and the peak position of sample B is 540 nm, while the main peak positionof sample C is 672 nm. In the case of sample C a smaller peak is visible in the extinction spectrum at 410 nm. We attribute these peaks to the dipolar resonances in the NPs, hence for sample
C a higher order mode is present at short wavelengths.20 It is seen that the peaks red-shift withincreasing average particle size, while the full width at half maximum (FWHM) increases. These
self-assembled Ag NPs have different sizes and spacing, and this will in principle give rise to
an inhomogeneous broadening of the LSP resonances from the distribution of the particle sizes.

FIG. 2. Extinction spectra for samples A, B and C. Reference PL spectrum is shown for comparison.

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Sample A has the most regular-shaped NPs with a small spread in size distribution, and is in effectshowing a resonance peak with the smallest FWHM. Resonance peak of sample C seems to belocated at around 700 nm, but has a large FWHM in contrast, the tail of which is covering the QWs
emission spectrum. Sample B had a better spectral overlap with the QWs emission than the other Ag
NPs structures. It is therefore expected that LSP coupling with QWs of sample B can be the mostpronounced. The resonance peaks shift towards longer wavelengths as Ag NP size increases dueto the red-shift of the dipolar resonance for larger diameter, as long as the dipole approximation isvalid.9
B. PL measurements

We investigated PL enhancement of GaInN/GaN QWs with various Ag NPs. For PL measurements the sample excitation and emission collection were done from the backside of the polishedsapphire substrate, and the spectra measured at room temperature are shown in Fig. 3. It is seen that
the PL emission of sample A has remained nearly unchanged compared to the bare LED emission,
while PL enhancements are observed for sample B and C. Increasing the Ag thickness, hence also
the average particle size, results in an increase in PL enhancement. The integrated PL enhancementsof samples A-C relative the reference sample are 1.25, 2.59 and 6.02, respectively.
Other mechanisms could exist behind the PL enhancement in addition to LSP-QW coupling.
One is an enhanced reflection from the top p-GaN surface of either the excitation laser signal,
and/or the emission from the QWs, both of which would give an enhanced signal at the detector.
The reflectance spectra (not shown) reveal that the Ag NPs formed on the p-GaN surface do not
result in an increased reflection compared to the as-grown surface at the QWs emission wavelength.
At the excitation wavelength the reflection is enhanced by 10 %, which results in an insignificant enhancement factor that is already taken into account in Fig. 3.20 These factors are therefore
excluded.
C. Temperature dependent PL

We proceed to study the improvement of the IQE based on temperature dependent PL. We
analyzed the IQE of the LEDs with various Ag NPs by measuring PL intensity versus temperature, as shown in Fig. 4. Under the assumption that non-radiative recombinations are inactivated
at cryogenic temperatures, one can retrieve the IQE from the temperature dependence of integrated PL intensity I.26 The IQE is defined by normalizing the integrated PL intensity to 1 at
20 K, i.e., assuming the IQE to be 100 %. Thus, the IQE at room temperature will be defined

FIG. 3. Room temperature PL intensity of GaInN/GaN 5QWs with Ag NPs on top. Excitation power density is 756 W/cm2,
and emission peak is around 530 nm.

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FIG. 4. Temperature dependence of integrated PL for estimating IQE. PL intensity has been normalized to 100 at the 20 K
value.

as IQE = I295K/I20K. As the temperature increases, PL intensity degrades due to the activation of
non-radiative recombination process. The IQE for the 530 nm emission was found to be 19.3 %for the reference sample. Regarding the Ag NP coated samples, the expression for the IQE will be
modified due to LSP-coupling, η = (krad + kSPη LSP)/(krad + knrad + kSP), where k rad, k nrad and kSP
are the radiative, non-radiative and LSP coupling rates, and η LSP is the radiative coupling efficiencyof the LSP mode. It has been demonstrated that kSP and knrad follow similar trends as a functionof temperature, and it is therefore reasonable to estimate k SP ∼ 0 at low temperatures and assume
that IQE is approximately 100 % for the Ag NP coated samples.27 We can then estimate the IQE ofsamples A-C the same way as the reference, and obtain 26.1 %, 26.4 % and 44.2 %, for samples
A, B and C, respectively, with 756 W/cm2 of excitation power density. Table I summarizes the IQEenhancement factors.
We note that the IQE of A is almost equal to that of B, and higher than the reference sample IQE, contrary to what was observed in the PL measurements of Fig. 3 where sample B had
a considerably higher PL intensity. The integrated PL intensity is proportional to the external
quantum efficiency, and as such it includes the effects of both IQE and light extraction efficiency(LEE). The observed difference between the integrated PL and IQE enhancement factor suggests
a modified LEE for samples A-C. The IQE enhancement of sample A is higher than its integrated
PL enhancement which indicates a degraded LEE. As for samples B and C is it seen that the integrated PL intensity has a higher enhancement than the IQE. Assuming the integrated PL intensity
is proportional to the external quantum efficiency (EQE), the LEE enhancement can be estimatedthrough the ratio between integrated PL enhancement and IQE enhancement. This follows since
the integrated PL enhancement can be written as IAg/Iref = LEEAg/LEEref × IQEAg/IQEref . The LEE
TABLE I. Summary of enhancement parameters. Also given is the LSP radiative efficiency and Purcell factor obtained frompower density variation and the relation between IQE enhancement and initial IQE.

SampleA
B
C

Avg. size[nm]Integrated PL
enhancementa

Deduced LEE
enhancement

High power IQE
enhancementb

Low power IQE
enhancementb

η LSP

FP

50
120
185

1.25
2.59
6.02

0.93
1.90
2.63

1.35
1.36
2.29

1.10
1.71
8.12

0.38
0.45

2.4
26.5

a At 295 K, and 756 W/cm2 excitation power.
bHigh power: 756 W/cm2, Low power: 1 W/cm2.

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ratios for samples A-C are listed in Table I, where sample B and C show an enhancement by a factorof 1.9 and 2.6, respectively, while A shows a reduced LEE based on these estimations.
The improved LEE due to Ag NPs is observed from sapphire-side emission, and is expected tobe coming from the scattering of Ag NPs acting on the QW emission, i.e. the radiatively emitted
light from the QWs that are not coupled to LSP modes. The emitted light from the QWs will see
a modified interface when Ag NPs are included compared to a bare GaN-air interface, and this is
expected to change the LEE.27 In the following we will direct our attention to the IQE enhancement
caused by the Ag NPs.
D. Excitation power density dependent IQE

We continue to study the excitation power density dependent IQE, which further supports theconclusion of improved IQE due to plasmonic coupling from temperature dependent PL measurement. In Fig. 5(a) we obtain the IQE at various excitation power densities for the emission pointwith 530 nm peak wavelength at 756 W/cm2. We first note the increasing IQE of the reference
sample with increasing power densities. This effect is attributed to coulomb screening of thequantum-confined stark effect (QCSE) by the increasing carrier density.28
Fig. 5(b) shows the IQE enhancements with changing excitation power density. The trend forsamples B and C is a decreasing enhancement with power density, while for sample A the opposite
trend is observed, although weak. It is also noticed that the enhancements factors for sample A andB approach each other at high power density.
To understand the decrease in IQE enhancement, we need to keep in mind the increase inoriginal (reference sample) IQE with increasing power density. The observed trend is that the lower
the initial IQE is, the higher the enhancement factor is due to LSP coupling, and this is in agreementwith the theoretical expectation of an inverse relation between IQE enhancement due to Ag NPs, K,
and initial IQE without Ag NPs, η i , which can obtained as the ratio,29
K=(krad + kSPη LSP) / (krad + kSP + k nrad) 1 + Fp η LSP=,
k rad/ (k rad + k nrad)
1 + Fp η i(1)where Fp = kSP/krad, is the Purcell enhancement factor, and η LSP is the radiative efficiency of the
LSP-mode. This relation also implies that the requirement for IQE enhancement, i.e. K > 1, theLSP-mode radiative efficiency has been to be larger than the initial IQE, η LSP > η i . We plot
the IQE enhancement factors of Fig. 5(b) as a function of the initial IQE in Fig. 6. The inverse
relation is most clear for sample C, while sample B shows a similar but weak trend. For sample A,
the opposite trend is observed. Using Eq. (1) we can roughly fit the data in Fig. 6 to obtain boththe Purcell factor and the radiative efficiency of the Ag NPs. From these measurements, we estimate
a Purcell enhancement factor of 26.5 and an LSP radiative efficiency of 45 % for the Ag NPs ofsample C. The values obtained for sample B are 1.7 for the Purcell and 38 % for LSP radiative

FIG. 5. (a) The estimated IQE and (b) IQE enhancement factor as a function of excitation power density at 295 K. Inset isthe zoom in for sample A and B.

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FIG. 6. IQE enhancement factor with initial IQE variation, calculated from the measurements of excitation power density
variation. Dashed lines are fittings based on Eq. (1). Inset is the zoom-in at small IQE enhancement factor for sample A and B .

efficiency. The results are summarized in Table I. Due to the opposite trend observed for sample A,
no parameters could be obtained. The difference in Purcell factors between sample B and C imply
different LSP-QW coupling rates, where C has a higher coupling rate than B. Through numericalcalculations we have previously investigated the decay rate enhancement in the plane of the QWsfrom Ag NPs, and have found that larger NPs exhibit larger enhancement factors for particle sizes
up to about 200 nm in diameter.30 The difference in average Ag NP size between sample B and C
explains why the Purcell factor of C is larger than B.
As seen from the extinction spectra, the LSP resonance of sample A is blue-shifted relative tothe emission peak, while those of B and C are red-shifted. When increasing the excitation power
density, in addition to the increase in IQE, the GaInN/GaN QWs peak emission wavelength isblue-shifted due to the screening piezoelectric field by free-carriers.31,32 For the reference sample
the emission peak around 528 nm is red-shifted to 537 nm when the excitation power is decreased to
1 mW/cm2. Such a red-shift with decreasing excitation power is also present for samples A-C. For
sample A this implies an emission peak shift towards the LSP resonance. As such, two competingmechanisms are present in the case of sample A when increasing the excitation power density; thefirst is the emission peak blue-shift, resulting in a better overlap with the LSP-mode resonance, andthe second is an increase in original IQE from the increased carrier density, resulting in a reduced
IQE enhancement factor from LSP-coupling. The increasing enhancement factor of sample A as
seen from the inset of Fig. 6 when excitation power density is increased can therefore be expected
to be the result of increasingly better spectral overlap between the QWs emission spectrum andthe LSP-mode. The extinction spectra for sample B and C are relatively broad (FWHM larger than
200 nm), hence this effect would be insignificant for B and C. However, the extinction FWHM of
sample A is around 50 nm, therefore an emission peak shift of about 10 nm is expected to have aneffect on the LSP-QW coupling.
We also note that despite having the best spectral overlap between sample B resonance peak
and the emission wavelength, sample B does not have the highest PL or IQE enhancement. Thiscan be understood by considering the absorption and scattering properties of the Ag NPs. These
effects were previously investigated and it was found that the scattering to absorption ratio is an
important characteristic when determining whether Ag NPs can provide efficiency enhancement.30
The absorption in the NP is the loss channel for the LSP mode, and reducing these losses relative
to the scattering can ensure a higher LSP mode radiative efficiency. Despite the matching of theextinction or resonance peak of sample B with the emission wavelength, sample C is expected
to have a higher scattering capacity at the emission wavelength due to its larger average Ag NP
sizes relative to sample B.30,33 This is because the scattering-to-absorption ratio of sample C ishigher than B due to the larger Ag NPs. As such sample C is scattering dominated and sample A
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into the LSP mode, the scattering efficiency of the NP is crucial when energy is to be re-emitted
into free-space radiation. This agrees with the parameters obtained by the fittings of Eq. (1), where
sample B and C had LSP radiative efficiencies of 38 % and 45 %, respectively.

IV. CONCLUSIONS

We investigated the LSP-QW coupling and resulting IQE enhancement of GaInN/GaN QW
LEDs with self-assembled Ag NPs. It was found that the strong PL enhancement was partly due
to LSP-QW coupling, and partly due to LEE enhancement, and separating these effects we noted
an IQE improvement due to LSP-QW coupling at 530 nm emission from 19.4% to 44.1% usinglarge sized Ag NPs (sample C) at 756 W/cm2. It was also found that the IQE enhancement isstrongly dependent on excitation power density, yielding highest enhancement factors at low free
carrier densities. Where an IQE enhancement by a factor of 2.3 was observed at 756 W/cm2,
an enhancement factor of 8.1 was observed at 1 W/cm2. Our results confirm the inverse relation
between the aforementioned quantities, and implicate that it is imperative to take the excitation
power density into account when conducting PL measurements to investigate the enhancement due
to LSP-coupling. We therefore establish quantitative enhancement factors depending on original
IQE, and further confirm the fact that the lower the IQE of a light emitting structure is, the higher
the enhancement factor will be due to surface plasmonics from Ag NPs.
ACKNOWLEDGMENTS

This work was partially supported by a Strategic Young Researcher Overseas Visits Programfor Accelerating Brain Circulation from Japan Society for the Promotion of Science, the Program
for the Strategic Research Foundation at Private Universities, 2012–2016, supported by the Ministryof Education, Culture, Sports, Science and Technology, and the Scandinavia-Japan Sasakawa Foundation. Y. Chen acknowledges financial support from the Danish Research Council for Technology
and Production Sciences (Grant No. 10093787), and the National Natural Science Foundation ofChina (Grant No. 61405067). A. Fadil and H. Ou acknowledge the support of Innovation Fund
Denmark (Grant No. 0603-00494B).
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Journal of Crystal Growth 351 (2012) 126–130

Contents lists available at SciVerse ScienceDirect

Journal of Crystal Growthjournal homepage: www.elsevier.com/locate/jcrysgro

MOVPE growth of nonpolar a-plane GaN with low oxygen contaminationand specular surface on a freestanding GaN substrate
Yasuhiro Isobe a, Hiromichi Ikki a, Tatsuyuki Sakakibara a, Motoaki Iwaya a,n, Tetsuya Takeuchi a,
Satoshi Kamiyama a, Isamu Akasaki a,c, Takayuki Sugiyama b, Hiroshi Amano b,c, Mamoru Imade d,
Yusuke Mori daFaculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japanc
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japand
Graduate School of Engineering, Osaka University, Osaka 565-0871, Japanba r t i c l e i n f o

a b s t r a c t

Article history:
Received 22 December 2011
Received in revised form
19 April 2012
Accepted 20 April 2012
Communicated by C. Caneau
Available online 10 May 2012

We investigated unintentionally doped nonpolar a- and m-plane GaN layers grown by metalorganic
vapor phase epitaxy under several sets of conditions on freestanding a- and m-plane GaN substrates.
Oxygen contamination in a-plane GaN is greatly reduced by increasing the V/III ratio during growth. As
a result, a high-resistivity GaN buffer layer for an AlGaN/GaN heterostructure ﬁeld-effect transistor was
realized.
& 2012 Elsevier B.V. All rights reserved.

Keywords:
A1. Impurities
A3. Metalorganic vapor phase epitaxy
B2. Semiconducting gallium compounds
B3. High electron mobility transistors

1. Introduction
Nonpolar group III nitride semiconductors have attracted
signiﬁcant attention for use in optical and electronic device
applications because of their lack of spontaneous and piezoelectric polarization ﬁelds [1]. Most III-nitride heterostructureﬁeld-effect transistor (HFET) devices are grown on c-plane [2–7]
and have two possible operation modes; depression mode
(D-mode) and enhancement mode (E-mode). D-mode HFETs are
already practically applied in high-frequency circuits, whereas
E-mode HFETs are expected to be used in future fail-safe circuitswith low power dissipation.
In the case of HFET devices in the nonpolar GaN, the twodimensional electron gas (2DEG) density in the channel layer can becontrolled by intentional doping in the barrier layers [8]. Therefore,
E-mode operation can be expected with a comparatively easy device
fabrication process. However, there are several problems in the useof a nonpolar plane for HFET devices. First, heteroepitaxially grown
nonpolar GaN ﬁlms have a high density of dislocations and stacking
faults [9,10]. This problem can be solved by performing growth on anCorresponding author.
E-mail address: iwaya@meijo-u.ac.jp (M. Iwaya).

0022-0248/$ - see front matter & 2012 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.jcrysgro.2012.04.030

freestanding GaN substrate [11]. Second, nonpolar planes are
chemically active. Therefore, contamination by oxygen is a serious
problem [12]. Oxygen acts as a shallow donor in GaN [13], making
the fabrication of HFETs with a high breakdown voltage difﬁcult[14]. Fe-doping in polar and nonpolar HFETs has been used tocompensate for the oxygen [15,16]. Although this is effective for
reducing the leakage current, it tends to degrade the device
performance through Fe segregation [17,18].
In this study, we investigated the residual oxygen contaminationin nonpolar GaN by varying the growth temperature and V/III ratio
during the growth of GaN by low-pressure metal organic vapor
phase epitaxy (MOVPE). In particular, we characterized the surface
morphology and oxygen contamination for each set of conditions.

2. Experimental details
All samples were grown by MOVPE. Trimethylgallium (TMGa),trimethylaluminum and ammonia (NH3) were used as the sourcegases. As the substrates, we used þc-, m- and a-plane freestanding
GaN ﬁlms grown by the Na-ﬂux method [19]. The threading
dislocation density and stacking fault density of these GaN substrates were determined to be less than 5  106 cm  2 and less than2  103 cm  1, respectively. These substrates were sliced with an

Y. Isobe et al. / Journal of Crystal Growth 351 (2012) 126–130

127

GaN substrates, respectively. After 10 min thermal annealing at
1030 1C in NH3 and hydrogen, an unintentionally doped GaN layer
was grown. In the growth of a- and m-plane GaN, the surface
morphology can be improved by performing growth under a low V/
III ratio of as low as 200 and at a low growth temperature of 1030 1C[22,23]. The growth conditions for a- and m-plane GaN were as
follows: V/III ratio: 250, growth temperature: 1030 1C, reactor
pressure: 100 hPa. Next, an unintentionally doped GaN layer was
grown at a V/III ratio of 1000, a growth temperature of 1030 1C and areactor pressure of 100 hPa. Samples 4 and 5 were grown on m- anda-plane Na-ﬂux GaN substrates, respectively. After 10 min thermal
annealing at 1060 1C in NH3 and hydrogen, an unintentionally doped
GaN layer was grown. The growth temperature and pressure were
ﬁxed at 1060 1C and 100 hPa, respectively. The V/III ratio wasincreased from 250 to 1000 and then 2000. Note that the growthrate on each plane is different under the same growth conditions
and that the total thickness is not exactly the same for all samples.
For example, the growth rate of sample 1 (þc-plane) was approximately 3 mm/h. In a growth temperature of 1030 1C (sample 2 andoffset angle of 070.51 and then subjected to chemical mechanicalpolishing. By using these freestanding GaN substrates, a GaN buffer
with a specular surface morphology and good crystalline quality canbe realized [20,21].

3. Results and discussion
Table 1 shows the growth conditions, the root mean square(RMS) surface roughness obtained by atomic force microscopy(AFM) and the oxygen concentration in the top GaN buffer
obtained by secondary ion microprobe mass spectrometry (SIMS).
Sample 1 was grown on a þc-plane GaN substrate. This sample wasprepared as a reference for comparison. After 10 min thermal
annealing at 1050 1C in NH3 and hydrogen, an unintentionally
doped GaN layer was grown. The growth conditions of GaN wereas follows: V/III ratio: 1000, growth temperature: 1050 1C, reactor
pressure: 200 hPa, TMGa ﬂow rate: 150 mmol/min, NH3 ﬂow rate of
3 slpm. Samples 2 and 3 were grown on m- and a-plane Na-ﬂux
Table 1
Growth conditions, surface roughness and oxygen concentration of epitaxial GaN.
Plane

Sample 1
Sample 2
Sample 3
Sample 4
Sample 5

þc
m
a
m
a

1st GaN layer

2nd GaN layer

3rd GaN layer

Temperature (1C)

V/III ratio

Temperature (1C)

V/III ratio

Temperature (1C)

V/III ratio

1050
1030
1030
1060
1060

1000
250
250
250
250

–
1030
1030
1060
1060

–
1000
1000
1000
1000

–
–
–
1060
1060

–
–
–
2000
2000

2nd layer

1st layer

Carbon concentration (atoms/cm3)

Oxygen concentration (atoms/cm3)

Epi-GaN of oxygen
concentration(each top layer)

0.29
0.16
0.19
0.82
0.24

Detection limit
2  1018
1  1017
3  1017
Detection limit

1E21

1E21
1E20
1E19
M-plane(sample 2)

1E18
A-plane (sample 3)

1E17
C-plane (sample 1)
detection limit

1E16

1E20
1E19
1E18
M-plane(sample 2)

1E17

A-plane (sample 3)
detection limit

1E16
1

0

2
Depth (µm)

3

4

0

1

2
Depth (µm)

3

4

1E21
rd

3

layer

2nd

Carbon concentration (atoms/cm3)

1E21
Oxygen concentration (atoms/cm3)

Surface roughness
5  5 mm2 RMS (nm)

st

layer

1 layer

1E20
1E19
M-plane(sample 4)

1E18
1E17

A-plane (sample 5)detection limit

1E16
0

1

2

3

5
4
6
Depth (µm)

7

8

9

1E20
1E19
1E18
M-plane(sample 4)

1E17
A-plane (sample 5)detection limit

1E16
10

0

1

2

3

5
4
6
Depth (µm)

7

8

9

10

Fig. 1. Impurity proﬁles of epitaxial GaN layers measured by SIMS. (a, c) and (b, d) show the oxygen and carbon concentrations in samples 1–5, respectively.

128

Y. Isobe et al. / Journal of Crystal Growth 351 (2012) 126–130

3), although the growth rate was slightly higher for m-plane, it wasalmost same as c-plane. However, the growth rate of m-plane wasapproximately 15 to 30% higher than that of a-plane under hightemperature growth condition.
Fig. 1 shows the oxygen and carbon concentration proﬁles ofeach epitaxial GaN layer in samples 1–5. The detection limit of theoxygen and carbon concentrations in SIMS measurement is
2  1016 cm  3. According to Fig. 1(a) and (b), the oxygen andcarbon concentrations in sample 1 ( þc-plane) were below thedetection limit. In contrast, samples 2 (m-plane) and 3 (a-plane)
exhibit a relatively high oxygen density. The oxygen concentrations of samples 2 (m-plane) and 3 (a-plane) are 2  1018 and1  1017 cm  3, respectively. Additionally, the carbon concentrations of samples 2 (m-plane) and 3 (a-plane) are 7  1017 cm  3and exceed the detection limit. Samples 1–3 are not signiﬁcantly
contaminated by carbon. The concentrations of silicon and hydrogen impurities are also below the detection limit at all samples.
To reduce the oxygen contamination, growth at a higher
temperature is effective [12,24–26]. However, when m- anda-plane GaN were grown at a growth temperature of over
1080 1C, the surface morphology deteriorated. Therefore, we ﬁxed
the growth temperature at 1060 1C and changed the V/III ratio tocontrol the oxygen concentration in samples 4 (m-plane) and 5(a-plane). To decrease the oxygen concentration with increase in
V/III ratio has already been reported [12], similarly, the SIMS
proﬁles of these samples show that it was suppressed under thesegrowth conditions. In particular, the top layer of sample 5 on
a-plane GaN (V/III ratio: 2000) has an oxygen concentration
below the detection limit. From these results, GaN is conﬁrmed
the difference in oxygen and carbon incorporation on a-, m- andc-plane GaN, differences in the incorporation of impurities withthe crystal orientation have also been reported for other materials[27]. Similarly, we consider that the surface kinetics varies withthe crystal orientation owing to the formation of dangling bonds.
Fig. 2 shows differential-interference-contrast microscopic(DIC) images and bird’s-eye-view AFM images of each sample.
Samples 1–3 exhibit specular surfaces in the DIC images. The AFM
results also indicate atomically ﬂat surfaces. The RMS surface
roughness values of þc-, m- and a-plane GaN were 0.29, 0.16 and0.19 nm, respectively. However, the surface morphologies ofsamples 4 and 5 were much rougher than those of samples 1–3.
Such rough surfaces would cause difﬁculties in device fabrication
processes.
Next, we grew a sample structure on a-plane GaN with thesimultaneous aim of reducing the oxygen concentration andrealizing a ﬂat surface. The growth rate has been reported toaffect the crystalline quality and surface morphology of c-plane
GaN [28,29]. Fig. 3 shows a schematic view of the sample

sample 1

structure. This sample structure was grown on a freestanding
a-plane GaN substrate by a newly developed V/III ratio modulation process. After 10 min thermal annealing at 1060 1C in NH3and hydrogen, an unintentionally doped GaN layer was grown.
The growth temperature and growth pressure were ﬁxed at1060 1C and 100 hPa, respectively. In contrast, the V/III ratio
was increased stepwise from 250 to 1000 and then 2000. The
thicknesses of the ﬁlms grown at each V/III ratio were set to 1,
1 and 2 mm, respectively. The growth rate was 3 mm/h. Then, a 1mm-thick GaN layer was grown at a low growth rate of approximately 1.0 mm/h and a high V/III ratio of 3500, where the growthrate was controlled by adjusting the TMGa supply. A thick andhighly resistive GaN buffer layer is indispensable even whenusing a GaN substrate owing to the n-type conductivity of thesubstrate and Si contamination at the interface between the GaN
substrate and the ﬁrst epilayer [30]. If the growth condition of the
low-growth-rate layer only use, the time of the device fabricationis very long. It is very important to fabricate the device in a shorttime is important. So, we combined the growth conditions ofconventional-growth-rate layer and the low-growth-rate layer.
Subsequently, a 1-nm-thick unintentionally doped Al0.36Ga0.64N
layer, a 3-nm-thick Si-doped Al0.36Ga0.64N layer and a 17-nmthick unintentionally doped Al0.36Ga0.64N layer were grown,

u-Al

Ga

N:17nm

n-Al

Ga

N:3nm

u-Al

Ga

N:1nm

Low growth rate layer>1um

V/III ratio:3500
Temperature:1060°C
Unintentionally doped GaN

V/III ratio:2000
Temperature:1060°C
Unintentionally doped GaN

V/III ratio:1000
Temperature:1060°C>4um

Unintentionally doped GaN

V/III ratio:250
Temperature:1060°C
Unintentionally doped GaN

A-plane LPE-GaN sub
Fig. 3. Sample structure on freestanding a-plane GaN substrate obtained using
newly developed V/III ratio modulation process.

sample 2

sample 3

sample 4

sample 5100µ m100µ m100µ m100µ m100µ msample 1

sample 2

sample 3

sample 4

sample 52nm

RMS=0.290nm

1.5nm

RMS=0.156nm

10nm

RMS=0.187nm

15nm

RMS=0.821nm

4nm

RMS=0.237nm

Fig. 2. Sample morphology of epitaxial GaN observed in (a) DIC images and (b) bird’s-eye-view AFM images.

Y. Isobe et al. / Journal of Crystal Growth 351 (2012) 126–130

129

5µ m × 5µ m

1.5nm

100µ mRMS=0.182nm

Counts/s

Concentration (Atoms/cm3)

1E21
Measurement
Fitting curve

1E20
1E19
Oxygen

1E18
Carbon1E17detection limit

1E16
55

56

57
58
59
2Theta/Omega

60

1

0

2

3
4
Depth (µm)

56Fig. 4. a-Plane epitaxial AlGaN/GaN structure in Fig. 3: (a) DIC microscopic image, (b) bird’s-eye-view AFM image, (c) high-resolution XRD spectrum with 2yXo-scandiffraction and (d) SIMS proﬁle of oxygen and carbon concentrations.

101<1-100><0001>

Au mask
AlGaN
GaN<11-20><1-100><0001>

10nm

Fig. 5. (a) Plan-view TEM and (b) cross-sectional TEM images of the structure in
Fig. 3.

where the V/III ratio, growth temperature and growth pressure
for each layer were 500, 1090 1C, and 100 hPa, respectively.
Fig. 4 shows the surface morphology of the sample shown in
Fig. 3 obtained by high resolution X-ray diffraction (XRD) with a(11–20) 2y/o-scan and SIMS proﬁle of the oxygen and carbon
concentrations of the sample. The RMS roughness of this sample
was 0.18 nm over an area of 5 mm  5 mm. A satellite peak was
observed in the XRD 2y/o-scan. As shown in Fig. 4(a) and (b), a
macroscopically specular surface and an atomically ﬂat surface
were obtained by inserting the low-growth-rate layer. Furthermore, the oxygen concentration was under the detection limit.
Fig. 5 shows plan-view and cross-sectional TEM images of thesample shown in Fig. 3. The threading dislocation and stacking
fault densities were determined to be less than 5  106 cm  2 andless than 2  103 cm  1, respectively. Thus, the crystalline qualityof epitaxial a-plane AlGaN/GaN is similar to that of the Na-ﬂux
GaN substrate.
Using this growth method, a high-performance nonpolara-plane GaN-based HFET can be realized on an unintentionally

IDS [A/mm]<11-20>500nm>1 10 17 cm-3 Vgs = -10 V(Standard growth condition)

100
10-1
10-2
10-3

<2 10 cm
(newly developed process)

10-4
10-5
0

10

20
VDS [V]30

40

Fig. 6. Off-breakdown voltage of the structure in Fig. 3 at a gate bias is  10 V. (Forinterpretation of the references to color in this ﬁgure, the reader is referred to the
web version of this article.)

doped GaN buffer layer. Next, we measured the off-breakdown
voltage of the structure in Fig. 3 at a gate bias of  10 V toestimate the speciﬁc resistance. The result is shown in Fig. 6.
Under the following device fabrication process conditions, mesa
isolation was performed by removing approximately 300 nm of
AlGaN and GaN by reactive ion etching using Cl2 gas. The ohmic
contact used for the drain/source, which was composed of Ti(30 nm)/Al (100 nm)/Ti (20 nm)/Au (150 nm), was annealed at
850 oC for 30 s in N2. The gate electrode was composed of Ni(20 nm)/Au (80 nm). The gate length, width and spacing between
the drain/source metals of the device were 2, 100 and 8 mm,

130

Y. Isobe et al. / Journal of Crystal Growth 351 (2012) 126–130

respectively. The blue line in Fig. 6 is the leakage current of thesample with an oxygen concentration of over 1  1017 cm  3,
which was grown under the standard growth conditions withthe growth conditions of the GaN buffer layer in sample 3. The redline is the leakage current of the sample grown by the newly
developed V/III ratio modulation process. These results show that
we successfully achieved devices with a very low leakage current
by using the new process. The resistivity of the newly fabricateda-plane GaN buffer layer is estimated to be higher than1  103 O cm. Details of other device performances have beenreported in a previous paper [8].

4. Conclusion
Control of the V/III ratio and the growth temperature iseffective for reducing oxygen contamination in an epitaxial GaN
ﬁlm as well as for realizing an atomically ﬂat surface on an
a-plane GaN substrate grown by the Na-ﬂux method. We wereable to fabricate a high-quality nonpolar a-plane AlGaN/GaN
wafer with a high-resistivity GaN buffer layer and a smoothsurface. The results of this study can be applied to the growthof high-resistivity GaN buffer layers in nonpolar a-plane HFETs.

Acknowledgments
This study was supported by New Energy and Industrial
Technology Development Organization.
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Journal of Crystal Growth 401 (2014) 367–371

Contents lists available at ScienceDirect

Journal of Crystal Growthjournal homepage: www.elsevier.com/locate/jcrysgro

Control of crystallinity of GaN grown on sapphire substrateby metalorganic vapor phase epitaxy using in situ X-ray diffraction
monitoring method
Motoaki Iwaya a,n, Taiji Yamamoto a, Daiki Tanaka a, Daisuke Iida a, Satoshi Kamiyama a,
Tetsuya Takeuchi a, Isamu Akasaki a,babFaculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

art ic l e i nf o

a b s t r a c t

Available online 2 December 2013

In this study, we investigated the application of a novel in situ X-ray diffraction (XRD) monitoring method
to control GaN crystal growth using a low-temperature (LT)-deposited buffer layer. We found that thismethod is useful in controlling the annealing of the LT-buffer layer, which strongly depends on the
crystallinity of GaN. Accordingly, if we employ in situ XRD grown on GaN using LT-buffer layer onsapphire substrate, the optimization of the annealing conditions will become easier because it would bepossible to determine by only one growth procedure. Therefore, we expect that this method will serve asa new and helpful optimization tool for crystal growth.
& 2013 Elsevier B.V. All rights reserved.

Keywords:
A1. Characterization
A1. X-ray diffraction
A3. Metalorganic vapor phase epitaxy
B1. Nitrides

1. Introduction
Nowadays, high-brightness blue, green, and white light-emitting
diodes (LEDs) [1] and high-power violet and blue laser diodes [2,3]based on group-III nitride semiconductors have been commercialized
as a result of several breakthroughs such as the growth of highcrystallinity GaN on sapphire using a low-temperature (LT)-depositedbuffer layer [4] and realization of conductivity control of nitrides[5–7]. Optoelectronics in the visible-short wavelength region has
been established via the use of nitrides. These devices are typicallyfabricated by metalorganic vapor phase epitaxy (MOVPE) [8].
In device manufacturing, a key process is an in situ monitoring
method for MOVPE. Such a method can be used to obtain information about some parameters during growth, such as the wafer bow,
crystalline quality, and lattice constant. This information provides
feedback about the growth condition and can be used to determine
the growth mechanism. Therefore, in situ monitoring is expected toserve as a key technology to improve the device performance.
To obtain this information during growth, in situ monitoring
methods based on different sources such as light, electron beams,
and X-rays can be used. Among these sources, an electron beam is notsuitable for MOVPE because it requires high vacuum. Thus, visible andinfrared light are typically used [9–13]. Thus far, many studies havefocused on in situ measurement methods using visible and infrarednCorresponding author.
E-mail address: iwaya@meijo-u.ac.jp (M. Iwaya).

0022-0248/$ - see front matter & 2013 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.jcrysgro.2013.11.010

light. Such methods can be used to measure the surface morphology,
layer thickness, wafer curvature, and surface temperature. For example, because the reﬂectance intensity oscillates with an increase in theﬁlm thickness, this interference effect allows us to determine thegrowth rate [14]. Because the reﬂectance intensity also signiﬁcantly
affects the surface morphology, it is effective for controlling thecoalescence of crystal nuclei during initial growth in the case of
GaN growth via an LT-buffer layer on a sapphire substrate [15]. It isalso possible to determine the critical ﬁlm thickness from curvature
measurement, such as an AlGaN/GaN heterostructure [16,17]. Moreover, in situ monitoring devices for MOVPE are commercially available,
and these systems have already been used to develop devices suchas LEDs.
However, visible and infrared light cannot be used for nanoscale monitoring. Therefore, methods based on them cannot be
used to measure the lattice constant, relaxation, and compositionand crystalline quality of ﬁlms. For example, when GaN ﬁlms weregrown on sapphire substrate using an LT-buffer layer, their crystallinity was found to be strongly dependent on the annealing
condition of the LT-buffer layer [18]. However, in situ monitoring
using visible and infrared light is not suitable for controlling theannealing condition of the LT-buffer layer.
In principle, in situ monitoring using X-rays is a candidate inthis regard. The resolution of X-rays differs signiﬁcantly from thatof light, and therefore, in situ monitoring based on the former canbe used to measure the lattice constant, thin-ﬁlm nanostructure,and crystalline quality. Several studies have already focused onin situ monitoring using X-rays for MOVPE growth [19–24].

368

M. Iwaya et al. / Journal of Crystal Growth 401 (2014) 367–371

However, most of these analyzed the crystal growth mechanism.
Studies have hardly focused on feedback about the crystal growth,such as crystallinity, via in situ monitoring using X-rays.
In this study, we investigated the application of a novel in situ
X-ray diffraction (XRD) monitoring method to control GaN crystal
growth using an LT-buffer layer. As a result, we found that thismethod is useful for controlling the annealing of the LT-buffer
layer, which strongly depends on the crystallinity of GaN.

2. Experimental procedure
GaN ﬁlms were grown on (0 0 0 1) sapphire substrates using
MOVPE apparatus with a horizontal water-cooled quartz reactor.
Trimethyl-gallium (TMGa) and ammonia (NH3) were used as Ga
and N source materials, respectively. Hydrogen (H2) and nitrogen(N2) were used as carrier gases. Table 1 and Fig. 1 show the GaN
growth conditions for MOVPE and the timing chart of the growth
process, respectively. After thermal cleaning of the sapphire
substrates in H2 gas at 1020 1C for 10 min, a  20-nm-thick GaN
LT-buffer layer was deposited at 520 1C for 3 min. The LT-buffer
layer on sapphire was heated up to 1125 1C for 6 min, and its
temperature was maintained with appropriate annealing time.
After annealing, a 3-μm-thick GaN ﬁlm was grown at 1125 1C for
80 min.
We evaluated the LT-buffer layer and GaN epitaxial layer withsymmetric (0 0 0 2) Bragg diffraction using an in situ XRD system.
We equipped the MOVPE system with an in situ XRD monitoring
device. A beryllium window was installed in the MOVPE to serve as a

Table 1
Growth conditions in this study.
Process

Reactor pressure [hPa]Gas(a) Thermal cleaning(b) Temperature down(c) GaN buffer layer

933
933
933

(d) Thermal annealing

933–500(e) GaN growth

500

H2
H2
TMGa
H2
NH3
H2
NH3
TMGa
H2
NH3

Fig. 1. Timing chart of GaN growth process in this study. (a) Thermal cleaning in H2
ﬂow at 1020 1C for 10 min. (b) Temperature down at 520 1C for 10 min. (c) LT-GaN
was deposited at 520 1C. (d) LT-GaN was annealed [Annealing time 1 (3, 6) min:Substrate was heated up to 11251 for 1 (3, 6) min. Annealing time 9 (12, 15) min:Substrate was heated up to 1125 1C for 6 min. Then, the temperature wasmaintained for 3 (6, 9) min. Therefore, the annealing time 9 (12, 15) is given by
the ramp time (6 min) plus the holding time (3 (6, 9) min).]. (e) 3-μm-thick GaN
was grown at 1125 1C.

viewport for passing X-rays. An in situ X-ray was focused on thesample surface using a Johansson curved crystal mirror. By thisapproach, the incidence angle of the X-ray can be changed without
moving the X-ray source. Moreover, the diffracted X-ray was
detected by a one-dimensional charge-coupled device [25,26]. By
this method, the scattered X-ray can be detected without moving
the substrate and detector. By using this conﬁguration, this system
realized the equivalent of a (0 0 0 2) 2θ/ω scan without requiring theuse of an analyzer crystal in 1 s during wafer rotation. Furthermore,
the 2θ value was calibrated by the XRD peak from the sapphire
substrate. In this system, the tilt component and distribution of thelattice constant c can be simultaneously characterized from the full
width at half maximum (FWHM) [23,24]. Although the resolution
limit decreases slightly, this in situ XRD system is excellent formeasuring the equivalent of a (0 0 0 2) 2θ/ω scan for a very shorttime and without moving the substrate, X-ray source, and detector.
In this study, we measured the in situ XRD spectrum at a frequency
of once per period. Further, FWHMs of the in situ XRD spectrum
were analyzed by using a Gaussian ﬁtting method. The growthtemperature was measured by in situ optical monitoring. We also
observed the crystallinity of GaN by a typical ex situ XRD
measurement.

3. Result and discussions
Fig. 2 shows the typical in situ XRD spectra with symmetric(0 0 0 2) diffraction of the LT-GaN buffer layers with different
annealing times (as deposited (0), 1, 3, 6, 9, 12, and 15 min). In
this study, we deﬁned an annealing time of 0 min (as deposition of
buffer layer) as the time at which buffer layer deposition iscompleted. These ﬁgures showed that the X-ray spectra of thedeposited buffer layer included a clear peak attributable to the
GaN crystal. This suggests the crystallization of hexagonal single
crystals in the LT-GaN buffer layers before annealing. According toprevious reports, these broadened in situ XRD spectra also suggestthat the crystalline structure of the LT-GaN buffer layers is
composed of a random amorphous-like phase, a cubic phase, ora mixed cubic-hexagonal phase [14,15,18,23,27–30]. Furthermore,
the in situ XRD spectrum of the annealed LT-GaN buffer layer
indicated that thermal annealing of the LT-GaN buffer layer
immediately after its growth serves to induce the crystallizationof hexagonal single crystals. Then, the intensity of the in situ XRD
spectrum was strongest when the annealing time was 6 min, after
which it began decreasing. In other words, the hexagonal single
crystals were evaporated by thermal annealing for an annealing
time exceeding 6 min. At this time, the crystal structure wasalmost completely transformed into hexagonal single-crystal
islands. Although these results have been reported previously[18], this in situ XRD method can be used to determine theoptimization of the growth condition through only one growth.
Fig. 3 and Table 2 summarize the peak intensity and FWHM of
the in situ XRD spectrum from LT-GaN buffer layers during thermal
annealing as a function of the annealing time. The ﬁgure showsthat the X-ray intensity increases up to an annealing time of 6 min,after which it decreases. In contrast, the FWHM narrows with anincrease in annealing time up to 9 min. The FWHM was theminimum at an annealing time of  9 min, and it saturated. The
FWHM then widened with a further increase in annealing time.
To consider the annealing effect of the LT-GaN buffer layer, wewere evaluated to prepared LT-GaN buffer layer samples withannealing time of 3, 6, and 12 min. Fig. 4 shows the scanningelectron microscopic (SEM) images of each sample. From theseﬁgure, we conﬁrmed a marked difference in the surface morphology by annealing. Fig. 5 shows the schematic view of
the changes due to annealing of the LT-GaN buffer layer by

1000

500

34

36

1000

500

0
32

38

1000

500

0
32

34

36

36

500

0
32

38

500

0
32

2θ[degree]34

36

2θ[degree]

34

36

38

2θ[degree]

1000

38

369

1000

2θ[degree]

X-ry intensity [counts/s]

X-ry intensity [counts/s]

2θ[degree]

34

X-ry intensity [counts/s]

032

X-ray intensity [counts/s]

X-ray intensity [counts/s]

XRD intensity [counts]

M. Iwaya et al. / Journal of Crystal Growth 401 (2014) 367–371

38

1000

500

0
32

34

36

38

2θ[degree]

Fig. 2. X-ray spectra from LT-GaN buffer layer with various annealing times by using in situ XRD monitoring method. (a) 0 min, (b) 3 min, (c) 6 min, (d) 9 min, (e) 12 min and(f) 15 min.

Fig. 3. Annealing time dependence of X-ray intensity and FWHM (0 0 2) fromin situ XRD monitoring method.

Table 2
Results of in situ and ex situ XRD.
Annealing
time
[min]

In situ XRD
intensity[counts]

LT-GaN
FWHM[arcs]

3
6
9
12
15

602
945
678
249
74

3560
2800
2800
2990
3080

FWHM [arcs]
Tilt[0 0 2]Twist[1 0 2]

506
285
256
280
266

823
314
310
325
368

in situ XRD and SEM results. We considered etching and crystallization in LT-GaN buffer layer is in progress simultaneously in theinitial time of annealing. And crystallization ends at a certain

period of time, while only etching continues. Accordingly, optimal
annealing time of LT-GaN buffer layer should be present in order toobtain a high-quality GaN epitaxial layer.
Next, we discuss the correlation between the crystallinity of the
GaN epitaxial layer and the in situ XRD measurement results. We
investigated the crystalline quality of 3-μm-thick GaN layers
formed on LT-GaN buffer layers with various annealing times. In
this investigation, the samples were prepared separately. Furthermore, for annealing times of 1 and 3 min, GaN ﬁlms were grownafter the temperature was increased to 1125 1C for 1 min after theend of annealing. Therefore, we thought that some error may exist
for annealing times of 1 and 3 min. Fig. 6 shows the annealing timedependence of the FWHM of ex situ X-ray rocking curves (XRCs)
from the (0 0 0 2) diffraction, which is a distribution of the tilt
component, and the (1 1 0 2) diffraction, which is a distributionof the twist component, of 3-μm-thick GaN ﬁlms. Although slighttilt component is also included from the (1 1 0 2) diffraction,
twist component is dominant in this measurement method[31,32]. Table 2 summarizes the FWHM of ex situ XRCs from the(0 0 0 2) and (1 1 0 2) diffractions as a function of annealing
time. The ﬁgure and table show that the tilt and twist distributionsof GaN ﬁlms are strongly dependent on the annealing time of the
LT-GaN buffer layer. The FWHMs of the ex situ XRC from GaN
narrowed with an increase in the annealing time up to 9 min. The
FWHM was the minimum at an annealing time of 9 min, and it
saturated. Then, the FWHM widened with a further increase in theannealing time. In other words, we found that the crystallinity of
GaN is the best when it was grown on the narrowest FWHM of theannealed LT-buffer layer. By using the same technique, we could
well reproduce the annealing time of the LT-buffer layer. Therefore, we found that this in situ XRD monitoring method is usefulfor controlling the annealing of the LT-buffer layer. Accordingly, if

370

M. Iwaya et al. / Journal of Crystal Growth 401 (2014) 367–371

3 μm

3 μm

3 μm

Fig. 4. Plan view SEM images of LT-buffer layer with annealing times of (a) 3 min, (b) 6 min, and (c) 12 min.

LT-buffer layer (nucleation)
Saphhire

Saphhire

Saphhire

Saphhire

Saphhire

Saphhire

0 min

3 min

6 min

9 min

12 min15 minEtching and crystallization
proceed simultaneously.

Crystallization is almost finished.
Etching continues.

Fig. 5. Schematic view of the changes due to annealing of the LT-GaN buffer layer from experimental results.

Acknowledgment
This study was supported by the Ministry of Education, Culture,
Sports, Science and Technology of Japan through the Supported
Program for the Strategic Research Foundation at Private Universities, 2012–2016.

References

Fig. 6. Annealing time dependence of FWHM of tilt (0002) and twist (1–102)
distributions from 3-μm-thick GaN ﬁlms on LT-buffer layers with different
annealing times.

we apply the in situ XRD monitoring system under various growth
conditions, the optimization of growth conditions will be easier,because it would be possible to determine the optimized growth
condition through only one growth. The in situ XRD monitoring
system proposed herein enabled us to estimate the optimized
annealing condition during growth. This implies that this system is
extremely effective for the optimization of crystal growth.

4. Conclusion
We examined the application of a novel in situ XRD monitoring
method to control GaN crystal growth using an LT-buffer layer. The
FWHM determined from the in situ XRD monitoring method could
be used to accurately observe the changes in the LT-buffer layer
with annealing. Accordingly, if we employ in situ XRD grown on
GaN using LT-buffer layer on sapphire substrate, the optimizationof the annealing conditions will become easier because it would
be possible to determine by only one growth procedure. We foundthat this in situ XRD monitoring method is useful for controllingthe annealing of the LT-buffer layer.

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Journal of Crystal Growth 393 (2014) 108–113

Contents lists available at ScienceDirect

Journal of Crystal Growthjournal homepage: www.elsevier.com/locate/jcrysgro

In situ X-ray diffraction monitoring of GaInN/GaN superlattice duringorganometalic vapor phase epitaxy growthTaiji Yamamoto a,n, Daisuke Iida a, Yasunari Kondo a, Mihoko Sowa a, Shinya Umeda a,Motoaki Iwaya a, Tetsuya Takeuchi a, Satoshi Kamiyama a, Isamu Akasaki a,babFaculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

art ic l e i nf o

a b s t r a c t

Available online 1 December 2013

We observed the growth of the Ga0.80In0.20N (2 nm)/GaN (3 nm) superlattice (SL) structure by in situ
X-ray diffraction (XRD) monitoring. The satellite peaks from the  1st to the þ 1st order can be obtained
from these in situ XRD spectrums. From the full width at half maximums (FWHMs) of the 0th and  1stsatellite peaks as a function of the SL periods, we observed a clear trend in each FWHM. It was found thatby analyzing this trend along with ﬂorescence microscopic and transmission electron microscopic
analysis, an analysis of the In segregation and misﬁt dislocation are possible. Accordingly, if we employ
in situ XRD under various growth conditions, the optimization of the growth conditions will become
easier because it would be possible to determine the number of periods at which In segregation andmisﬁt dislocation increases by only one growth procedure.
& 2014 Published by Elsevier B.V.

Keywords:
A1. X-ray diffraction
A3. Organometallic vapor phase epitaxy
B1. Nitrides
B3. Heterojunction semiconductor devices

1. Introduction
Because the bandgap of GaInN ternary alloys has a broad range
from 0.65 eV to 3.43 eV [1,2], these alloys are suitable for solar cellapplications. By using a combination of GaInN alloys, one cantheoretically design a multijunction solar cell with a bandgap
spanning the infrared to ultraviolet regions of the solar spectrum[3–5]. This makes it possible to achieve high-efﬁciency multijunction solar cells. Thus far, by improving the crystallinity of
GaInN, improving the electrode structure, and using the concentration of sunlight, the conversion efﬁciency of GaInN-based solar
cells has been improved by up to 4% [6]. We have reported that thesuperlattice (SL) structure is useful for improving the crystallinityof the GaInN active layer [7–12]. However, there are numerous
practical problems that must ﬁrst be solved. Although an increasein the SL periods and InN molar fraction is required for realizing
high-performance GaInN-based solar cells, the crystallinity of
GaInN SL will deteriorate rapidly by introduction of misﬁt dislocations [3], introduction of growth pits [13,14], and increase of In
segregation [13–17] in GaInN, if the layer exceed the critical
thickness. In particular, In segregation (composition ﬂuctuation)
will rapidly degrade the device performance such as GaInNbased blue and green laser diode and solar cells [16,17]. Also,
In segregation can be measured from the dot emission densitynCorresponding author. Tel.: þ 81 528321151.
E-mail address: 133434039@ccalumni.meijo-u.ac.jp (T. Yamamoto).

0022-0248/$ - see front matter & 2014 Published by Elsevier B.V.
http://dx.doi.org/10.1016/j.jcrysgro.2013.11.072

determined by ﬂuorescence microscopy. We found that the solar
cell characteristics degrade rapidly when this dot emission density
exceeds  106 cm  2 [17]. Therefore, it is very important todetermine the optimal growth conditions under which the dot
emission density is less than 106 cm  2. In addition to the dot
density, the device characteristics are also degraded by thepresence of misﬁt dislocations. Therefore, numerous experiments
will be required to explore the optimal growth conditions unlessin situ observations are employed.
In situ monitoring in organometallic vapor phase epitaxy(OMVPE) is a key process in device manufacturing. This approach
can provide information about the growth condition and help
clarify its mechanism. Therefore, in situ monitoring is expected toserve as a key technique for improving the device performance.
Thus far, in situ observation methods using visible/infrared lighthave been used to improve the device performances [18–21].
In situ monitoring can be used to measure the surface morphology,
layer thickness, wafer curvature, and surface temperature. However, it is impossible to perform nanoscale measurements using
visible/infrared light, because its wavelength exceeds several
hundred nanometers. Therefore, we cannot measure with thistechnique the crystalline quality and structure of thin ﬁlm layersuch as SLs structures. In contrast, because the resolution of X-ray
and visible/infrared light is signiﬁcantly different, in situ monitoring using X-rays can be used to measure the SL structure. Several
studies have reported on in situ observations using X-rays in
OMVPE growth [22–27]. However, most of these have focused on
the analysis of the crystal growth mechanism.

T. Yamamoto et al. / Journal of Crystal Growth 393 (2014) 108–113

2. Experimental procedure
The samples were grown on c-plane sapphire substrates using
OMVPE apparatus with a horizontal face down 2″  3″ reactor(TNEMC: GRC-203). Trimethyl-gallium (TMGa), trimethyl-indium(TMIn), and ammonia (NH3) were used as Ga, In, and N source
materials, respectively. Hydrogen (H2) and nitrogen (N2) wereused as the carrier gases. Revolution speed of the substratesusceptor was 5 rpm in constant. After the growth of a 3-μmthick GaN template at 1050 1C on c-plane sapphire covered with alow-temperature GaN buffer layer [28,29] using H2 carrier gas at
933 hPa, it was cooled to 750 1C and the carrier gas was changed toN2. Then, 75-period Ga0.80In0.20N (2 nm)/GaN (3 nm) SL was grownon the GaN template. The growth time of Ga0.80In0.20N and GaN
are 42 s and 81 s, respectively. During growth of the Ga0.80In0.20N
layer, the ﬂow rates of the TMG of 65 μmol/min, TMI of 65 μmol/
min, NH3 of 800 mmol/min, V/III ratio of 8500, and N2 of 16 l/min
were used. The threading dislocation density of the GaN template
was 3  108 cm  2. We evaluated the GaInN ﬁlms with symmetric (0002) Bragg diffraction using an in situ XRD system [26].
Fig. 1 shows the schematic view of the experimental setup.
We attached in situ XRD observation equipment to the OMVPE
apparatus. The X-ray monitoring system used in this study was

supplied by Rigaku Corporation. The OMVPE apparatus was
equipped with a beryllium window that served as a viewport for
passing X-rays. An in situ X-ray was focused on the sample surface
using a Johansson curved crystal mirror. By using this method, theincidence angle of the X-rays can be changed without moving the
X-ray source. Moreover, the diffracted X-ray was detected by a
one-dimensional charge-coupled device. By using this method, thescattered X-ray can be detected without moving the substrate anddetector. By using this conﬁguration, this system realized the105

broadening
0th

104

Intensity (cps)

In this study, we have applied in situ X-ray diffraction (XRD)
measurement to optimize the GaInN SL structure via the suppression of In segregation and the generation of misﬁt dislocations.

109

103

-1st
10SLs
18SLs
23SLs
29SLs
33SLs
43SLs

103
101
100

32

33

34

2θ (degree)
Fig. 3. Enlarged view of Fig. 1 at 0 and  1satellite peak region.

Total GaInN thickness (nm)Substrate
spread X-ray

Slit

Cu X-ray tube

20 30 40 50 60 70 80

diffracted X-ray

0th
-1st

0.34

Johansson curved crystal
focus the X-ray beam on
the growing sample.

FWHM(degree)

OMVPE reactor

X-ray 1D CCD detector

Substrate

0.32
0.30
0.28
0.26

Focused X-Ray

0.24

10 15 20 25 30 35 40
Periodic number

Contain many
angle X-rays

Fig. 4. FWHM of the 0th and  1st satellite peaks as a function of the SL periods.

Diffracted X-Ray

Total GaInN thickness (nm)Fig. 1. Schematic view of the experimental setup.

20 30 40 50 60 70 80

FWHM(degree)

0.30
0.25
0.20
0.15

0th
-1st
simulation (0th)0.10
10 15 20 25 30 35 40
Fig. 2. In situ XRD spectrum of 75-period Ga0.80In0.20N (2 nm)/GaN (3 nm) SL
during growth of 10, 23, 33, and 43 periods.

Periodic number
Fig. 5. This ﬁgure was inserted the results of the simulation in Fig. 4.

110

T. Yamamoto et al. / Journal of Crystal Growth 393 (2014) 108–113

equivalent of a (0002) 2θ/ω scan without requiring the use of ananalyzer crystal in 1 s during the rotation of the wafers. Furthermore, the 2θ value was calibrated based on the GaN peak.
Although the resolution limit decreases slightly, this in situ XRD
system is able to perform the equivalent of a (0002) 2θ/ω scan at aresolution of 1arcsec. In this setup, the full half width maximum(FWHM) is both controlled by the dispersion in the lattice constantc and the mosaicity of the crystal [26,27]. Due to the diffraction
geometry, we measure here the mosaicity tilt. In this study weacquired in situ XRD data at the frequency of one diffractogram per
SL period. Further, FWHMs of the in situ XRD spectrum wereanalyzed by using a Gaussian ﬁtting method.

10 μm

10 μm

10 μm

10 μm

10 μm

Fig. 6. Fluorescence microscopic images of (a) 10, (b) 20, (c) 25, (d) 30, and (e) 40-period GaInN SL samples.

0.5×106 cm-2

0.9××106 cm-2

10 μm

10 μm

2.1×106 cm-2

1.0×106 cm-2

10 μm

2.1×106 cm-2

10 μm

10 μm

Fig. 7. Images corresponding to those shown in Fig. 5 (a)–(e) with blue light removed by using image processing. We have also shown the cyan dot emission density in eachﬁgure. (a) 10 peroids, (b) 20 periods, (c) 25 periods, (d) 30 periods and (e) 40 periods. (For interpretation of the references to color in this ﬁgure legend, the reader is referredto the web version of this article.)

T. Yamamoto et al. / Journal of Crystal Growth 393 (2014) 108–113

We also fabricated additional ﬁve Ga0.80In0.20N (2 nm)/GaN(3 nm) SL samples with 10,20,25,30, and 40 periods for ex situ
characterization. We evaluated the In segregation in each sample
through ﬂuorescence microscope (Nikon LV100D, Excitation light:
Halogen lamp (50 W), excitation power density: several kW/cm2)
observations. And we used transmission electron microscopy (TEM)(HITACHI H-9000A) to observe misﬁt dislocations. The accelerating
voltage of our TEM is 300 kV. In this condition, plan-view TEM
images were observed.

3. Results and discussion

Dots emission density (cm-2)

Fig. 2 shows the in situ XRD spectrum obtained from 75-period
Ga0.80In0.20N (2 nm)/GaN (3 nm) SL during the growth of 10, 23,
33, and 43 periods. Also, Fig. 3 shows an enlarged view of Fig. 2 at
0 and 1 satellite peak region. We conﬁrmed clear satellite peaks
from the  1st to the þ1st order. Even though the measurement
time was 1 s, the satellite peaks from the SL were observed andseparated. Therefore, reasonably high resolution and short measurement time was achieved by this in situ XRD system. In
addition, the FWHMs of the satellite peaks showed a certain
trend. Up to a certain period, the FWHM of the satellite peaks in
the in situ XRD spectrum reduced. Because the FWHM of XRD
generally decreases with increasing ﬁlm thickness, this is areasonable result in theory [30]. In contrast, FWHM increased
after a certain period. To exactly analyze this trend, we analyzed
the FWHMs of the satellite peak of each period. Fig. 4 summarizes
the FWHM of the 0th and  1st satellite peaks as a function of theSL periods. Moreover, Fig. 5 was inserted the results of the

2.0x106

1.5x106

1.0x106

5.0x105
10 15 20 25 30 35 40

Periodic number
Fig. 8. Dot emission density as a function of the SL periods.

111

simulation (FWHM of 0th satellite peak) in Fig. 4. In this simulation, we have used the XRD simulator (Rigaku: GlobalFit). There isa difference between simulation and experimental results. We
thought this difference is the effect of crystallinity of the underlying GaN layer and disorder of the GaInN/GaN superlattice.
In particular, it is considered that the inﬂuence of crystallinity isgreater because this in situ XRD does not use an analyzer crystal.
However, the trend of the experimental results and the simulationis the almost same up to a certain period. We observed a clear
trend in each FWHM. Up to 25 periods, the FWHMs from satellite
peaks of XRD decreased continuously with an increase in the
period. This trend is the same as the simulation result. However,
we conﬁrmed that the FWHM did not show this tendency andincreased after  25 periods. Therefore, we considered that acrystallographic change may have occurred. Next, we analyzed
the additional SL samples having different periods. Throughvarious analyses, we concluded that the FWHM increases with
the In segregation. Fig. 6 show ﬂuorescence microscopic images of
10-, 20-, 25-, 30-, and 40-period GaInN SL structures. These ﬁgures
show signiﬁcant differences. To clarify these differences, weapplied image processing techniques to these ﬂuorescence microscopic images. Fig. 7 show the images corresponding to thoseshown in Fig. 6 with the main blue emission from the SL removed
by image processing. These ﬁgures showed cyan emission attributable to In segregation. The evolution of the dot emission
density is summarized Fig. 8. It can be seen that the cyan dot
density strongly increases after 30 SL periods. This particular SL
thickness correlates with the onset of the FWHM increase
observed by XRD during the in situ XRD measurement (Fig. 4).
The main blue emission from the SL follows an opposite variation
with a strong decrease for SL thicknesses over 30 periods. Thiseffect is attributed to the generation of misﬁt dislocations, as
observed later by TEM (Fig. 9). Fig. 10 show ex situ XRD reciprocal
space maps of the asymmetric reﬂection (20–24) of sample with25, 30 and 40 SL periods. The SL relaxation is not directly apparent
with this ex-situ technique. In addition, we performed TEM
analysis in order to analyze in detailed microstructure. Fig. 9 show
plan-view TEM images at 25 and 30 period SL samples. Note that
we could observe to the interface of GaN and GaInN, because thethickness of the TEM sample is approximately 150 nm. In 25 SLs
sample, threading dislocations from the underlying GaN can be
conﬁrmed, but there was little misﬁt dislocations at GaInN/GaN
interface. In contrast, in 30 SLs sample, many misﬁt dislocations at
GaInN/GaN interface were conﬁrmed. These experimental results
could be explained as follows. Fig. 11 show the diagram of thegrowth model based on the obtained results. Up to 25 periods,
few dot emissions attributable to In segregation are observed.
In addition, misﬁt dislocations are not increased. Thus, the FWHM

Fig. 9. Plan vies TEM images of 25 and 30 period SL samples.

112

T. Yamamoto et al. / Journal of Crystal Growth 393 (2014) 108–113

Fig. 10. XRD reciprocal space mapping images from 25, 30 and 40 period SL samples.

only one growth procedure. The in situ XRD monitoring system
proposed herein enabled us to estimate the In segregation andmisﬁt dislocations of the GaInN/GaN SL structure during growth.
This implies that this in situ XRD observation is extremely effective
for the optimization of the device performance.

High

Up to 25 periods
Increased the period

In segregation
rapidly increase

Misfitdislocations
generated
Exceed 25 periods

Acknowledgments
This study was supported by the Ministry of Education, Culture,
Sports, Science and Technology of Japan through the Supported
Program for the Strategic Research Foundation at Private Universities, 2012–2016.

Fig. 11. Growth model diagram based on the obtained results.

from satellite peaks reduces in accordance with the increase inperiod, whereas In segregation and misﬁt dislocations rapidly
increases after 25 periods. Because the periodicity of the SL isdisturbed, the FWHM of the satellite peak increases. Accordingly,
if we employ in situ XRD under various growth conditions, theoptimization of the growth conditions will become easier because
it would be possible to determine the number of periods at which
In segregation and misﬁt dislocation increases by only one growth
procedure. In addition, this result conﬁrmed the improvement in
the device performance. The device performance is reported indetail elsewhere [31].

4. Summary
We observed the growth of the Ga0.80In0.20N (2 nm)/GaN(3 nm) SL by in situ XRD monitoring. The satellite peaks from the1st to the þ 1st order can be obtained from these in situ XRD
spectra. From the FWHMs of the 0th and  1st satellite peaks as afunction of the SL periods, we observed a clear trend in each
FWHM. Up to 25 periods, the FWHMs from the satellite peaks
decreased with an increase in the period. We can theoreticallydescribe this result. However, we conﬁrmed that the FWHM did
not continue to show this tendency and that it increased after  25
periods. This has been attributed to In segregation and misﬁt
dislocations based on ﬂuorescent microscopic and TEM analysis.
Accordingly, if we employ in situ XRD under various growth
conditions, the optimization of the growth conditions will become
easier because it would be possible to determine the number of
periods at which In segregation and misﬁt dislocations increase by

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physicaPhys. Status Solidi C 9, No. 3–4, 753–756 (2012) / DOI 10.1002/pssc.201100491

cwww.pss-c.com

current topics in solid state physics

Laser lift-off of AlN/sapphirefor UV light-emitting diodes

Hiroki Aoshima*,1, Kenichiro Takeda1, Kosuke Takehara1, Shun Ito1, Mikiko Mori1, Motoaki Iwaya1,
Tetsuya Takeuchi1, Satoshi Kamiyama1, Isamu Akasaki1,3, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Department of Electrical Engineering and Computer Science, Nagoya University, Nagoya 464-8603, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received 1 August 2011, revised 12 October 2011, accepted 31 October 2011
Published online 14 February 2012
Keywords UV LED, laser lift-off* Corresponding author: e-mail 103434001@ccalumni.meijo-u.ac.jp

We report on laser lift-off (LLO) of AlN/sapphire for UV
light-emitting diodes (LEDs). Underfill between chip andsubmount is a key factor for the successful LLO of
AlN/sapphire. We fabricated thin-film-flip-chip UV
LEDs with a peak wavelength of 343 nm using the LLO.

Moreover surface texturing was carried out on the exfoliated AlN surface. The light output power from the UV
LED after the LLO and the surface texture at exfoliatedsurface is 1.7 times higher than that before the LLO at 20mA.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction
UV light-emitting diodes (LEDs) are promising forvarious applications, such as bioagent detection, water andair purification, and dermatology [1]. Group III-nitridebased UV LEDs have been reported by several groups [26]. Although the internal quantum efficiencies (IQEs) of
DUV/UV LEDs are higher than 70% by improving the
crystalline quality of AlGaN [7], the external quantum efficiency (EQE) is still low. This is one of the most serious
problems preventing the realization of high-efficiency devices. Therefore, technology for the improvement of thelight extraction efficiency (LEE) is very important for
high-efficiency DUV/UV-LEDs.
At present, there are many reports on improving theLEE in nitride-based LED by the laser lift-off (LLO)
method [8], many of which are GaN/sapphire LLO, and the
LLO of AlN/sapphire was few [9,10]. Moreover, details of
the LLO conditions such as wavelength, irradiation condition, and implementation method have not been reported
yet. In addition, LLO of AlN is considered to be more difficult than that of GaN, because the melting point of AlN ismuch higher and precipitated aluminum by LLO of AlN isdifficult to remove.
In this study, we investigated LLO of AlN/sapphire indetail. We established the substrate separation of
AlN/sapphire just by laser irradiation as a simple process.
We also fabricated thin-film-flip-chip UV LEDs with apeak wavelength of 343 nm by LLO. Moreover surface
texturing was carried out at the exfoliated surface.
2 Experiments
UV LEDs were fabricated on a double side polishedsapphire (0001) substrate by metal organic vapor phase
epitaxy. First, we grew 2-μm-thick AlN on a sapphire substrate at 1400 °C, then, a 2.5-μm-thick Al0.25GaN0.75 film
was grown at 1100 °C. Next, the wafer was removed fromthe reactor, and grooves along the (1 100) direction wereformed by conventional photolithography and reactive ion
etching. The width, spacing, and depth of the grooves were5, 5, and 1 μm, respectively. Then, a thick Al0.25Ga0.75N
layer was regrown for 4 h on the grooved underlying layer.
The growth temperature and pressure during Al0.25Ga0.75N
growth were 1100 °C and 6.6×104 Pa, respectively. During
the growth of the AlGaN layer, the flow rates of TMGa,
TMAl and ammonia were 0.030, 0.013, and 61 mmol/min,
respectively. UV-LEDs were fabricated on this highcrystalline-quality AlGaN underlying layer. Figure 1
shows a schematic of the device structure in this study. The
thickness of the Si-doped n-Al0.25Ga0.75N layer was 4 μm.
Then, an unintentionally doped GaN (3 nm)/Al0.15Ga0.85N:Si (8 nm) three quantum well (QW) active
layer, a p-Al0.5Ga0.5N (20 nm) blocking layer, a pAl0.25Ga0.75N (50 nm) cladding layer, and a p-GaN (50 nm)
contact layer were successively stacked. The size of the© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidi754

H. Aoshima et al.: Laser lift-off of AlN/sapphire for UV light-emitting diodes

LED was 500 × 500 μm2 with a p-contact area of 500 ×
140 μm2. After exposing the n-Al0.2Ga0.8N contact layer to
Cl2 RIE, a Ti/Al/Ti/Au (30/100/20/150 nm) contact was
deposited on the exposed n-Al0.25Ga0.75N contact layer. The
n-electrode annealing conditions were 800 °C in nitrogenambient for 30 s. A Ni/Au (10/40 nm) contact pad was deposited on the p-GaN contact layer. The p-electrode annealing conditions were 550 °C in O2 ambient for 1 min.
The emission peak wavelength of the UV LED was 343
nm.
After dividing the above wafer into chips using scribing equipment, these chips were mounted with the flip-chip
configuration using gold bumps. We prepared two samples
with and without underfill as shown in Fig. 2. The underfill(U8437-2) was used to fill the gap between a device and a
submount. Here, the major component of the underfill is
epoxy resin. To remove the sapphire substrate, the LLO
technique was used. An ArF pulsed excimer laser with awavelength of 193 nm and pulse width of 15 ns was directed through the back of the transparent sapphire substrate. The laser light was absorbed at the interface between the AlN and the sapphire, and it will induce the decomposition of the AlN. The laser beam spot size was 1 ×2 mm2, which was slightly larger than the size of the flipchip UV LED. Figure 3 shows the irradiated energy density vs yield rate. The yield rate was obtained from the result of experiments by changing the irradiated energy density per ten samples each. The best yield rate was obtainedin the range where the irradiated energy density was higher
than 1.0 J/cm2. In this study, we employed the irradiatedenergy density was 1.0 J/cm2, which was sufficiently highto perform LLO with one laser pulse. The devices wereseparated between the substrate and device only by the irradiation of the laser, unlike in the case of GaN/sapphire.
Moreover, the radiation energy required to perform LLO of
GaN has been reported that approximately 0.26 J/cm2 [11].
This difference in energy density are probably due to thedecomposition temperature of AlN and GaN.
After the LLO, surface texturing will be carried out on
the backside surface by hot aqueous potassium hydroxide(KOH) wet etching for the improvement of LEE. The wet
etching condition was 100 °C for 1 min.

gold bumpsapphire
deviceSubmount

underfillsapphire
devicesample[b]sample[a]Figure 2 Flip-chip configuration with and without underfill.
100

80

Yield rate [%]statusphysica

pss

c

60

40

20

0
0.00

0.25

0.50

0.75

1.00

1.25

1.50

1.75

2.00

2

Irradiated energy density [J/cm ]Figure 3 Irradiated energy density vs yield rate.

3 Results and discussion
Figures 4(a) and 4(b) show photographs of the sample
after the LLO with and without the underfill, respectively.
These results clearly show a significant difference between
the samples with and without the underfill. While the sample without underfill was unintentionally chipped, thesample with underfill was successfully peeled off the sapphire substrate. We suspect that the physical shock by decomposition causes chip damage in the case without underfill. On the other hand, in the case with underfill, thedevice had no mechanical damage, because the UV LED
had been completely fixed to the mounting substrate,
Therefore, the underfill is useful to successfully performthe LLO.

With underfill

Without underfill

After LLO

Original chip size

500mm

p-GaN contact layer

p-Al0.25Ga0.75N
p-AlGaN electron blocking layer

GaN/AlGaN MQW
n-Al0.25Ga0.75N

Al0.25Ga0.75N

AlN
Sapphire(0001)

500mm

Figure 4 Photograph of UV-LED after LLO [left] with and[right] without underfill.

Figure 1 Structure of UV LED.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

www.pss-c.com

Contributed
ArticlePhys. Status Solidi C 9, No. 3–4 (2012)755[a] before LLO

Intensity(normalized)

1.0

Intensity [arb.unit]

Sapphire

0.8

AlN

AlGaN

before LLO
after LLO

0.6[b] after LLO

0.4
0.2
300

30

32

34

36

38

40

42

400
500
Wavelength[nm]

600

2θ/ω

Figure 7 Normalized EL spectra before and after the LLO.
Figure 5 (a) and (b) show 2θ-ω scan data from XRD before LLO
and after LLO.

To confirm the backside materials before and after LLO,
we measured X-ray diffraction (XRD) of the sample backside surfaces. Figures 5(a) and 5(b) show 2θ/ω scan datafrom XRD before and after LLO, respectively. A strong
peak from the sapphire substrate was observed in Fig. 5(a).
In contrast, after LLO, no peak from the sapphire substrate
was observed and intense AlGaN and AlN peaks were observed. This XRD result indicates no sapphire was remained on the backside after the LLO.
Figure 6 shows an SEM image of an AlN surface of oursamples after hot KOH wet etching at 100 °C. A large triangular texture was obtained. This texture structure is expected to improve the LLE in UV-LED.

Figure 8 shows I-V and L-I curves of UV LEDs before
the LLO, after the LLO, and after the LLO with KOH wet
etching under DC bias condition. While the operating voltages are almost the same, the light output power after the
LLO is 1.3 times higher than that before the LLO This improvement probably comes from the surface roughness after the LLO, as shown in the SEM image (Fig. 9). As a result, we did not clearly observe the damage in the performances of the UV LED. Also, the light output power after
the LLO with KOH wet etching is 1.7 times higher thanthat before the LLO at 20 mA current injection. This result
indicates that the LLE was improved and no damage wasintroduced by LLO and KOH etching.

6

1.00

Voltage(V)

0.75

6μm

Figure 6 SEM image of exfoliated AlN plane after KOH etching.

Figure 7 shows normalized electroluminescense (EL)spectra of UV LEDs before and after the LLO. The EL
emission peaks from each device were 343 nm. The peak
position and FWHM of the EL spectra are almost the same
before and after the LLO. This result indicates that serious
damage to the active layer was not introduced by the LLO.

www.pss-c.com

4
0.50
2

0

0.25

0

10
Current (mA)

Light Output(arb.unit)

before LLO
after LLO
after LLO with KOH

0.00
20

Figure 8 V-I and L-I curves of UV LEDs before LLO, after
LLO, and after the LLO with KOH wet etching.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss

c

756

H. Aoshima et al.: Laser lift-off of AlN/sapphire for UV light-emitting diodes

[3]

6μm

Figure 9 Surface SEM image of exfoliated AlN plane after the
LLO.

4 Conclusions
We have achieved the LLO of AlN/sapphire and the surface texture fabricated by wet etching for higher extraction
efficiency. Underfill is a key to successfully implement the
LLO of AlN/sapphire. The light extraction efficiency of
the UV-LEDs was also improved by 1.7 times by a combination of LLO and KOH etching.

H. Tsuzuki, F. Mori, K. Takeda, M. Iwaya, S. Kamiyama, H.
Amano, I. Akasaki, H. Yoshida, M. Kuwabara, Y. Yamashita, and H. Kan, J. Cryst. Growth 311, 2860 (2009).
[4] Y. Taniyasu, M. Kasu, and T. Makimoto, Nature 441, 325(2006).
[5] D. Morita, M. Sano, M. Yamamoto, T. Murayama, S. Nagahama, and T. Mukai, Jpn. J. Appl. Phys. 41, L 1434
(2002).
[6] A. Khan, K. Balakrishnan, and T. Katona, Nature Photon. 2,
77 (2008).
[7] K. Ban, J. Yamamoto, K. Takeda, K. Ide, M. Iwaya, T. Takeuchi, S Kamiyama, I. Akasaki, and H. Amano, Appl.
Phys. Express 4, 052101 (2011).
[8] T. Fujii, Y. Gao, R. Sharma, E. L. Hu, S. P. DenBaars, and
S. Nakamura, Appl. Phys. Lett. 84, 855 (2004).
[9] S. Hwang, D. Morgan, A. Kesler, M. Lachab, B. Zhang, A.
Heidari, H. Nazir, I. Ahmad, J. Dion, Q. Fareed, V. Adivarahan, M. Islam, and A. Khan, Appl. Phys. Express 4,032102 (2011).
[10] V. Adivarahan, A. Heidari, B. Zhang, Q. Fareed, M. Islam,
S. Hwang, K. Balakrishnan, and A. Khan, Appl. Phys. Express 2, 092102 (2009).
[11] T. Ueda, M. Ishida, S. Tamura, Y. Fujimoto, M. Yuri, T.
Saito, and D. Ueda, Phys. Status Solidi C 0(7), 2219 (2003).

References[1][2]

H. Hirayama, J. Appl. Phys. 97, 091101 (2005).
C. Pernot, M. Kim, S. Fukahori, T. Inazu, T. Fujita, Y. Nagasawa, A. Hirano, M. Ippommatsu, M. Iwaya, S. Kamiyama, I. Akasaki, and H. Amano, Appl. Phys. Express 3,061004(2010).

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

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Study on Efficiency Component Estimation of 405 nm Light Emitting Diodes from
Electroluminescence and Photoluminescence Intensities

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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JL16
http://dx.doi.org/10.7567/JJAP.52.08JL16

Study on Efficiency Component Estimation of 405 nm Light Emitting Diodes
from Electroluminescence and Photoluminescence Intensities
Kazuki Aoyama1 , Atsushi Suzuki2 , Tsukasa Kitano2 , Satoshi Kamiyama1 ,Tetsuya Takeuchi1 , Motoaki Iwaya1 , and Isamu Akasaki1;31Faculty of Science and Technology, Meijo University, Nagoya 468-0073, Japan
El-seed, Nagoya 468-0073, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received October 24, 2012; revised March 21, 2013; accepted March 22, 2013; published online July 22, 2013
The Shockley–Read–Hall (SRH) model was applied to the determination of photoluminescence (PL) and electroluminescence (EL)characteristics. From the ratio of the internal quantum efficiency (IQE) obtained from the PL and EL intensities, the carrier injection efficiencies(CIE) for 405 nm LEDs were derived. All the efficiency components including the IQE, CIE, and light extraction efficiency for 405 nm LEDs wereobtained for various structural parameters by fitting the experimental data to theoretical equations of the SRH model.
# 2013 The Japan Society of Applied Physics

1. Introduction

The external quantum eﬃciency (EQE) of LEDs is themost commonly used and important value, because it can beeasily measured. However, it is a product of the internal
quantum eﬃciency (IQE), carrier injection eﬃciency (CIE),
and light extraction eﬃciency (LEE). Therefore, each ofthem should be properly clariﬁed to understand the physics ofdevice operation. For the estimation of IQE, the determination of temperature-dependent PL intensity1) and the application of the Shockley–Read–Hall (SRH) model to theevaluation of EL intensity versus current characteristic2) havebeen reported. In the former approach, the determination of
PL excitation similarly to that in LED operation is diﬃcult.
For the latter case, the obtained IQE is a type of nominal
IQE including CIE. Therefore, the components of the actual
IQE and CIE cannot be derived. Estimation of CIE as theratio of the IQE for electroluminescence (EL) to that forphotoluminescence (PL) obtained using the SRH model has
been reported for an AlGaN/GaN ultraviolet laser diode.3)
However, the CIE for visible LEDs has not been investigated because it is not thought to be a signiﬁcant issue. For
the optimization of device structures, understanding of the
CIE and other eﬃciency components will be very useful.
In this study, we applied the SRH model to determinationof PL and EL characteristics, and derived the IQE, CIE, andLEE in 405 nm LEDs for various AlN molar fractions in the
p-AlGaN electron-blocking layer.

where J is the injected current density into the device, q is
the electron charge, and d is the total width of the quantum
well (QW). Optical excitation is used for the determinationof PL characteristic, where the rate of generation (G) by
photo excitation is expressed by5)
G ¼ An þ Bn2 :The third component in Eq. (1) is neglected, because there isno carrier overﬂow in photo excitation. The light intensities
of EL and PL are deﬁned as

J¼ An þ Bn2 þ Cn3 ;
qd

ð3Þ

2

ð4Þ

where kc and  are the sensitivity and a proportional
constant, respectively. Substituting Eqs. (3) and (4) into
Eqs. (1) and (2), respectively, the following equations are
obtained:
J ¼ L1=2 þ L þ L3=2 ;pﬃﬃﬃﬃﬃﬃﬃ
G ¼ P1 IPL þ P2 IPL ;

ð5aÞ
ð6aÞ

where
qdA
qd
; ¼
;
1=2
kc
ðkc BÞ
A
1
P1 ¼ pﬃﬃﬃﬃﬃﬃ ; P2 ¼ :

B

¼

¼

qdC
;
ðkc BÞ3=2

ð5bÞ
ð6bÞ

In the case of experimental PL, a nominal generation rate G
is expressed by
G¼

ð1Þ

L ¼ kc B;
IPL ¼ Bn ;2. Estimation of CIE by SRH ModelIn the SRH model, the nonradiative recombination factor A,
radiative recombination factor B, and droop factor C are
included as functions of the carrier density n. C is the droop
factor including the carrier overﬂow and auger recombination. However, it is known that the droop is negligibly small
in optimized 405 nm LEDs;4) therefore, the droop is thought
to be only caused by the carrier overﬂow in unoptimized
heterostructures in 405 nm devices. In the case of EL,
electron excitation is calculated from the current injection as
follows:2)

ð2Þ

Plaser ð1  RÞl Plaser ð1  RÞ
¼;
Aspot lhAspot hð7Þ

where Plaser is the excitation power density, R is the Fresnel
reﬂectivity,  is the absorption coeﬃcient, l is the total widthof the QW, and h is the photon energy of the excitation
source. These parameters such as , , , P1 , and P2 , wereobtained by ﬁtting of theoretical and experimental curves ofthe output versus carrier density characteristic for EL andthe light output versus generation rate characteristic for PL.
The absorption coeﬃcients for GaN and InN were assumed
to be 2:0  105 cm1 (Ref. 6) and 1:5  105 cm1 (Ref. 7),
respectively. That of the GaInN QW active layer was
1:29  105 cm2 , obtained from Vegard’s law using thevalues of the binary materials. Since the excitation source

08JL16-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JL16

K. Aoyama et al.

was a He–Cd laser with a wavelength of 325 nm, carriers
were generated not only in the QWs but also in the wholelayer, and carriers ﬂow into QWs from neighboring layers.
Therefore, Eq. (7) could not be used for the estimation of net
G and related parameters. However, we used the nominal G
to draw the curves of IPL as a function of G, and ﬁt Eq. (6a)to the curve to obtain the parameters P1 and P2 , which arealso the nominal values. In this ﬁtting, these curves arealmost identical in the lower carrier density range because of
the small contribution of the C component in Eq. (1). Then,the IQEs for EL and PL excitation were calculated bysubstitution of the parameters as described above into
Bn2
L
¼ ;
2
2
J
An þ Bn þ Cn
Bn2
Bn2
IPL
IQEPL ¼¼ P2:¼
An þ Bn2GG

IQEEL ¼

ð8Þ
ð9ÞFig. 1. (Color online) Samples structure.

Although G and P2 are nominal values, as mentioned
above, the ratio of P2 to G in Eq. (9) does not include theuncertainty of the excitation. The ratio of IQEEL to IQEPL
represents CIE. LEE was calculated as the ratio of theexperimental EQE to IQEEL .

Light output (mW)1

3. Experimental Procedure

NiAu

measurement

0.8

NiAu

theoretical

0.7

ITO measurement

0.6

ITO theoretical

0.5
0.4
0.3
0.2
0.1
0
0

1

2
Current density (A/cm-2)

3

4

(a)
180000160000Light output (arb.unit)

Firstly, the estimation of eﬃciency components for two
types of 405 nm LED with diﬀerent LEEs was examined.
From the same LED epitaxial wafer, an LED with Ni/Au
as a p-type electrode and an LED with ITO8) as a p-type
electrode were fabricated.
Secondly, 405 nm LEDs with an ITO electrode anddiﬀerent AlN molar fractions of 0, 15, and 25% in the
p-AlGaN electron-blocking layer (EBL) were fabricated,9)and the eﬃciency components of the devices were
examined. All the LED samples have a multiple quantum
well (MQW) active layer composed of ﬁve QWs, where thewidths of the QW and barrier layer are ﬁxed at 2.8 and12 nm, respectively. In addition, the samples with AlN molar
fractions of 15 and 25% have an EBL with a thickness of
30 nm.
The LED chips were mounted on TO headers without
resin encapsulation, as shown in Fig. 1, and their EL
properties were measured using an integrated sphere. Their
PL properties were measured using the excitation sourceof a 325 nm He–Cd laser and a cooled Si-CCD-type
photodetector.

0.9140000measurement

120000theoretical

100000
80000
60000
40000
20000
0
5E+22

5E+23

5E+24

Generation rate (cm-3s-1)(b)

4. Results and Discussion

Figures 2(a) and 2(b) show the curve-ﬁtting results ofEL and PL measurements for LEDs with ITO and Ni/Au
p-electrodes, respectively. The theoretical curves obtainedwith Eqs. (5a) and (5b) can be made to closely ﬁt theexperimental ones by adjusting the ﬁtting parameters , ,, P1 , and P2 . The modiﬁed carrier density in the PL
measurement is determined by ﬁtting the curve of IQEPL
versus carrier density to the curve of IQEEL versus carrier
density in the low-excitation range as shown in Fig. 3.
Also, Fig. 3 shows the ﬁtting results of samples with AlN
molar fractions of 0, 15, and 25%. From these results, theeﬃciency components when the carrier density was injected
at 50 mA in this sample structure were obtained, as shownin Table I. The LEE of the ITO device is about 15%, which

Fig. 2. (Color online) Fitting results for (a) EL and (b) PL.

is much higher than that of the Ni/Au device. This isreasonable because ITO has a very low light absorption. The
values of IQEEL and CIE for the two types of LED wereconsistent because they originated from the same epitaxial
wafer. Accordingly, these eﬃciency components are thought
to be reasonable, indicating the usefulness of our estimation
model.
For the LEDs with diﬀerent AlN molar fractions, the sameprocedure as above was carried out. Figure 4 shows the CIEsof the three LEDs as a function of carrier density. For thedevices with AlN molar fractions of 15 and 25%, the CIE at

08JL16-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JL16

K. Aoyama et al.
Table II. Comparison of eﬃciency components for LEDs with AlN molar
fractions of 0, 15, and 25% in p-AlGaN electron-blocking layer (injection
current is 50 mA).

10
9

PL IQE

8EL (Al:0%) IQE7EL (Al:15%) IQE

IQE (%)6

EL (Al:25%) IQE

AlN 0%

AlN 15%

EQE

3.7

9.1

AlN 25%
6.1

IQEEL

23.3

57.1

39.6

5

IQEPL

94.8

94.8

94.8

4

LEE
CIE

15.9
30.1

15.9
62.6

15.4
45.0

3
2
1
0
1.00E+14

1.00E+15

1.00E+16

Carrier density (cm-3)

Fig. 3. (Color online) Fitting results for IQEEL and IQEPL .

Table I. Comparison of EQE, IQECIE, and LEE for LEDs with ITOelectrode and NiAu electrode as p-electrode (injection current is 50 mA).

EQE
IQEEL

ITO device

Ni/Au device

9.1
57.1

7.1
56.3

IQEPL

94.8

94.8

LEE

15.9

12.5

CIE

62.6

61.0

100
Al:0%
Al:15%
Al:25%

90
80

CIE (%)

70
60
50
40
30

50 mA in this sample structure are summarized in Table II.
Although the CIE is strongly dependent on the AlN molar
fraction in the p-AlGaN electron-blocking layer, the LEEs inthis series are almost the same. This result is also reasonable.
Our estimation model for the eﬃciency components appears
to be reliable for various types of LED.
5. Conclusions

The SRH model was applied to 405 nm LEDs. The CIE
was derived from the ratio of the two types of IQE for the
EL and PL intensities. These eﬃciency components were
estimated for two series of LEDs. The eﬃciency componentsof LEDs with Ni/Au and ITO electrodes as p-electrodes
were compared. It was found that the LEE of the LED with
the ITO electrode was higher than that of the device with
the Ni/Au electrode, while IQEEL and CIE for the two
types of LED were consistent. The eﬃciency componentsof LEDs with AlN molar fractions of 0, 15, and 25% in
the p-AlGaN electron-blocking layer were also compared.
Although CIE is strongly dependent on the AlN molar
fraction in the p-AlGaN electron-blocking layer. The LEEs
in this series are almost the same. These results arereasonable. Therefore, derivation of the eﬃciency components by this approach is considered to be useful forunderstanding device physics.
Acknowledgment

20
10
0
2.00E+14

2.00E+15

2.00E+16

2.00E+17

2.00E+18

Carrier density (cm-3)This study was supported by Ministry of Education, Culture,
Sports, Science and Technology Supported Program for theStrategic Research Foundation at Private Universities, 2012–2016.

Fig. 4. (Color online) Carrier density dependence of CIE.

a low carrier density was almost 100% and decreased withincreasing carrier density. The CIE of the LED with an AlN
molar fraction of 15% slowly decreases with increasing
carrier density. However, the LED with an AlN molar
fraction of 0% exhibits a very low CIE even at low carrier
densities. This may contain some measurement error. The
PL measurement was carried out for only one sample, since
the active layer is assumed to be uniform over the 2-in. LED
wafer. However, IQEEL and IQEPL were not thought to beagreed with each other, indicating that ﬁtting errors may be
involved. At high carrier densities, the CIE of this device isconsidered to be reliable. All the eﬃciency components inthis series of LEDs when the carrier density was injected at

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# 2013 The Japan Society of Applied Physics

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Room-temperature CW operation of a nitride-based vertical-cavity surface-emitting laser using
thick GaInN quantum wells

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REGULAR PAPER

Japanese Journal of Applied Physics 55, 05FJ11 (2016)
http://doi.org/10.7567/JJAP.55.05FJ11

Room-temperature CW operation of a nitride-based vertical-cavity surface-emitting
laser using thick GaInN quantum wells
Takashi Furuta1*, Kenjo Matsui1, Kosuke Horikawa1, Kazuki Ikeyama1, Yugo Kozuka1, Shotaro Yoshida1,
Takanobu Akagi1, Tetsuya Takeuchi1, Satoshi Kamiyama1, Motoaki Iwaya1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering and Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

2

Received November 5, 2015; accepted January 22, 2016; published online April 19, 2016
We demonstrated a room-temperature (RT) continuous-wave (CW) operation of a GaN-based vertical-cavity surface-emitting laser (VCSEL) using
a thick GaInN quantum well (QW) active region and an AlInN/GaN distributed Bragg reﬂector. We ﬁrst investigated the following two characteristics
of a 6 nm GaInN 5 QWs active region in light-emitting diode (LED) structures. The light output power at a high current density (>10 kA/cm2) fromthe 6 nm GaInN 5 QWs was the same or even higher than that from standard 3 nm 5 QWs. In addition, we found that hole injection into the farthest
QW from a p-layer was sufﬁcient. We then demonstrated a GaN-based VCSEL with the 6 nm 5 QWs, resulting in the optical conﬁnement factor of
3.5%. The threshold current density under CW operation at RT was 7.5 kA/cm2 with a narrow (0.4 nm) emission spectrum of 413.5 nm peakwavelength. © 2016 The Japan Society of Applied Physics

1.

Introduction

III–nitrides and their alloys with direct bandgaps cover abroad range of operating wavelengths. After the ﬁrst
demonstration of nitride-based pn-junction LEDs in 1989,
GaInN-based LEDs and lasers were successively achieved.1–3) Now, GaInN-based edge-emitting lasers are operatingfrom ultraviolet to green regions.4–7) On the other hand,vertical-cavity surface-emitting lasers (VCSELs) were invented by Soda et al. in 1979, and the advantages of the
VCSELs are low power consumption and two-dimensional
arrays.8) Thus far, the power conversion eﬃciency of infrared
VCSELs reached 62%.9) Recently, nitride-based VCSELs
have also been achieved,10–16) and expected to be utilized
in retinal scanning displays, adaptive laser headlights, andvisible light communication systems. However, further
improvements of nitride-based VCSELs are required forpractical use. The nitride-based VCSELs still have a muchhigher threshold current density than the infrared VCSELs.17)
One of the possible reasons for the high threshold current
density is the low optical conﬁnement factor caused bythin quantum wells (QWs) in the nitride-based VCSELs.
Generally, in infrared VCSELs, an 8 nm 3 QW active regionis used within a 1λ-cavity, resulting in a more than 3% opticalconﬁnement factor.17) On the other hand, in nitride-based
light-emitting devices, a thin (less than 3 nm) GaInN QW
active region has been typically used to prevent a small
overlap between electron and hole wave functions due topiezoelectric polarization ﬁelds in the wells.18) If the active
region with 3 nm GaInN 5 QWs is used in a 4λ-cavity, theoptical conﬁnement factor is calculated to be only 2%. Notethat such a large cavity length should be used for the nitridebased VCSELs since an intracavity contact structure must be
used owing to the absence of conductive distributed Bragg
reﬂectors (DBRs), and this situation aggravates the low
optical conﬁnement factor. The use of thick GaInN QWscertainly results in the improvement of the optical conﬁnement factor, but also leads to a couple of issues related tocarrier injections. One of the issues is the small overlap
between electron and hole wave functions in the GaInN
QWs as mentioned above. Another is poor carrier injections, especially hole injections, into the farthest QW from ap-layer.19–22) Nitride-based materials have largely asymmetriccarrier properties, such as eﬀective mass and mobility,
between electrons and holes; thus, the issue is more serious
than in the case of the other III–V semiconductor materials.
In this study, we ﬁrst calculated the GaInN well width
dependence on the optical conﬁnement factor in the nitridebased VCSELs. Then, we experimentally measured the light
output power of 6 nm GaInN 5 QWs in an LED structure incomparison with that of 3 nm 5 QWs to investigate the eﬀectof thick QWs on the light output power, in other words, theoverlap between electron and hole wave functions in the
QWs. We also prepared an LED with speciﬁc 6 nm GaInN 5
QWs in which 4 QWs were emitting 405 nm wavelengthand the farthest QW from a p-layer was emitting 425 nm
wavelength. This is to determine whether the hole injection
into the farthest QW was suﬃcient by comparing the 425 and405 nm intensities from the LED. The device characteristic
measurements were carried out at a very high current density
to investigate the situation under a possible VCSEL operation. Finally, we demonstrated a room-temperature (RT)continuous-wave (CW) operation of a nitride-based VCSEL
containing 6 nm GaInN 5 QWs and a 4λ-cavity with an
AlInN=GaN DBR.
2.

Experimental methods

Figure 1 shows calculated optical conﬁnement factors as afunction of GaInN well width. A VCSEL structure in thecalculation consisted of a GaN substrate, a 40-pair AlInN=
GaN DBR, a 4λ-cavity, and an 8-pair SiO2=Nb2O5 DBR. The
4λ-cavity contained a 300 nm n-GaN layer, GaInN=6 nm
GaN 5 QWs, a 15 nm p-AlGaN electron-blocking (EB) layer,
a 60 nm p-GaN layer, a 10 nm p++-GaN contact, a 20 nm
indium–tin oxide (ITO) p-contact, and a 30 nm Nb2O5 spacer.
The thicknesses of the layers were designed for a wavelengthof 410 nm. Note that this structure is identical to that of a
VCSEL fabricated and measured later. We calculated theoptical power distribution in the cavity and DBRs, and thenestimated the ratio of optical power in all the QWs to that in
the cavity=DBRs. A 2.0% optical conﬁnement factor was
estimated in the case of 3 nm 5 QWs, while a 3.5% opticalconﬁnement factor was obtained in the case of 6 nm 5 QWs.
Thus, the calculation result suggests that 6 nm QWs have a

05FJ11-1© 2016 The Japan Society of Applied Physics

T. Furuta et al.
SiO /Nb O DBR
2

5

SiO2

4
3

5

Nb2O5 spacer
Cr/Ni/Aup-GaN

EB layer

4λ

2
1

n-GaN

4
5
Well width [nm]GaN/AlInN DBR

6

GaN substrate

Fig. 1. Calculated optical conﬁnement factors of a nitride-based VCSEL
as a function of well width.

high optical conﬁnement factor, which is comparable to thatof infrared VCSELs (∼4%).17) Thus, we decided to use 6 nm
5 QWs in the following experiments.
Next, we fabricated two LED test structures (Samples A
and B) in order to investigate the eﬀect of thick well width
on light output power, in other words, an overlap between
electron and hole wave functions. In this study, the LED test
structures were grown on c-plane sapphire substrates bymetal–organic chemical vapor deposition (MOCVD). The
test structure consisted of a 30 nm low-temperature-deposited
buﬀer layer, a 2 µm undoped GaN layer, a 0.4 µm n-type GaN
layer, a GaInN=GaN 5 QW active region, a 15 nm p-type
AlGaN EB layer, a 60 nm p-type GaN layer, and a 10 nm
p++-GaN contact. Sample A contained 6 nm GaInN=6 nm
GaN 5 QWs and Sample B contained 3 nm GaInN=6 nm
GaN 5 QWs. The peak wavelengths of both the 5 QWs were
adjusted to be 405 nm. After the epitaxial growth, LEDs
were fabricated with the wafers by our standard fabrication
process. The emission area of the LEDs was 15 µm in
diameter for the LED characteristic measurements under highcurrent densities, which is comparable to that in a possible
VCSEL operation. A transparent 20 nm ITO p-electrode wasformed on top of the p-GaN contact layer, while a Cr=Ni=Au
multilayer was used as both an n-electrode and a p-pad
electrode. Standard 300 µm2 LEDs were also fabricated formeasurement under low current densities.
Another LED test structure (Sample C) with speciﬁc 6 nm
GaInN 5 QWs was also prepared.19–23) The speciﬁc GaInN 5
QWs consisted of 4 QWs emitting a 405 nm peak wavelengthand the remaining QW emitting a 425 nm peak wavelength,
which is located at the farthest position from the p-AlGaN EB
layer (nearest to the n-GaN layer) in order to evaluate hole
injection into the farthest QW by monitoring an emission
intensity of not only 405 nm but also 425 nm. The two
diﬀerently sized LEDs mentioned above were also prepared.
Finally, we fabricated VCSELs with 6 nm GaInN 5 QWs
as shown in Fig. 2. The VCSEL structure was grown on a
GaN substrate by MOCVD, in the same manner as that
described in the optical conﬁnement factor calculation. Notethat the 40-pair AlInN=GaN DBR showed 99.7% reﬂectivity
at around 410 nm as measured with an absolute reﬂectivity
measurement system. The size of the VCSEL was deﬁned
by an 8-µm-diameter SiO2 aperture. A transparent 20 nm ITO
p-electrode was formed on the SiO2 aperture. In addition, aNb2O5 spacer layer and an 8-pair SiO2=Nb2O5 DBR were
deposited on the ITO p-electrode to form the VCSEL cavity.

Fig. 2. (Color online) Schematic view of our VCSEL with 6 nm 5 QWs
and an AlInN=GaN DBR.

Light output power [arb.unit]

0
3

2

ITO

5
4

6 nm QWs

3

3 nm QWs

2
1
0

0

2

4

6

8

2

Current density [kA/cm ]

10

(a)

Light output power [arb.unit]

Optical confinement factor [%]Jpn. J. Appl. Phys. 55, 05FJ11 (2016)1612

3 nm QWs

846 nm QWs
0
0.00

0.04

0.08

2

Current density [kA/cm ]

0.12

(b)
Fig. 3. (Color online) J–V–L characteristics of Samples A (6 nm 5 QWs)
and B (3 nm 5 QWs) at (a) high and (b) at low current densities.

The reﬂectivity of the 8-pair SiO2=Nb2O5 DBR was
measured to be over 99.9% at around 410 nm.
3.

Results and discussion

The current density–light output power (J–L) characteristics
of Samples A (6 nm QWs) and B (3 nm QWs) are shown in
Fig. 3(a) under high current densities up to 10 kA=cm2 and(b) under low current densities up to 100 A=cm2. The J–L
characteristics under high current densities clearly show thatthe light output power from the 6 nm QWs is comparable toor even higher than that from the 3 nm QWs. At the sametime, the light output power from the 3 nm QWs was muchhigher than that from the 6 nm QWs under low current
densities as typically observed in standard LED operations.
Thus, we conclude that the overlap between electron and hole
wave functions in the 6 nm QWs is suﬃciently high under

05FJ11-2© 2016 The Japan Society of Applied Physics

T. Furuta et al.
5026Voltage [V]Intensity [arb.unit]

10 kA/cm

5

30

4

20
2
0

1

40

10
0

5

10

2

0
15

Light output power [̭W]

Jpn. J. Appl. Phys. 55, 05FJ11 (2016)

Current density [kA/cm ]0.1

Fig. 6. (Color online) J–V–L characteristics of a VCSEL with 6 nm QWs
and an AlInN=GaN DBR under RT CW operation. The threshold current
density was 7.5 kA=cm2.

0.02
400
420
Wavelength [nm]440

Intensity [arb.unit]

380

Fig. 4. (Color online) Emission spectra of Sample C (speciﬁc 6 nm
5 QWs) under various current densities. Emission spectra of Sample A(uniform 6 nm 5 QWs) and emission spectra only from the farthest QW in
Sample C were also plotted.

410

0.4
Intensity ratio

I=1.1 Ith
I=1.0 Ith
I=0.9 Ith
I=0.8 Ith

0.3

415

Wavelength [nm]420Fig. 7. (Color online) Emission spectra before and after the threshold.
After the threshold, a narrow peak with 0.4 nm FWHM at 413.5 nm wasobserved.

0.2
0.1
0.0

0.01
0.1
1 2 10Current density [kA/cm ]Fig. 5. Ratio of emission intensity from the farthest QW to total emission
intensity from speciﬁc 5 QWs as a function of current density.

high current densities, like in the VCSEL operating range. It
seems that the screening of the piezoelectric polarization ﬁeld
by the injected carriers is suﬃcient under such high current
densities.
Figure 4 shows emission spectra of Sample C (speciﬁc
6 nm 5 QWs) under various current densities. In addition,
emission spectra of Sample A, in which all the 6 nm QWs
were identical and emitting 405 nm peak wavelength, areplotted with black dotted lines for comparison. Here, theintensities of all the spectra were normalized in their peakintensities. It is possible to estimate the emission spectra only
from the farthest QW emitting 425 nm peak wavelength inSample C by subtracting the spectra of Sample A from thoseof Sample C. Then, the spectra only from the farthest QW are
also plotted with colored lines. Figure 5 then shows the ratioof the integral intensity of the emission spectra from thefarthest QW (shown with colored lines in Fig. 4) to theintegral intensity of the emission spectra from Sample C(shown with black solid lines in Fig. 4) as a function ofcurrent density. As intensity ratio of more than 35% was
observed from 50 A=cm2 to 5 kA=cm2. In principle, a 20%
intensity ratio is expected in the case of a uniform carrier
injection into 5 QWs, and the measured value here washigher than the expected value. We consider that the higher
value could come from a smaller bandgap of the farthest QW,
which should cause a larger carrier capture. In addition, we
observed that the ratio was relatively low when the current
density was less than 50 A=cm2 and more than 5 kA=cm2.
The low carrier injections into the farthest QW should be
due to low hole injections. Such low hole injections may becaused by the existence of large piezoelectric ﬁelds at the
low current density resulting from the small screening of theﬁelds and by carrier overﬂow, in other words, a poor hole
injection, at the high current density. Eventually, we
conclude that the carrier injection into the 6 nm 5 QWs is
suﬃciently uniform.
Figure 6 shows the J–V–L characteristics of the VCSEL
with 6 nm GaInN 5 QWs under CW operation at RT. The
threshold current density is 7.5 kA=cm2, and the operating
voltage at the threshold is 4.8 V. The diﬀerential quantum
eﬃciency (ηd ) is still low, 0.46%, with the maximum output
power of 45 µW at 14 kA=cm2. At this moment, themaximum output power could be limited owing to the selfheating of the VCSEL, considering the low thermal conductivity (∼5 W m−1 K−1) of AlInN layers.24) Figure 7 showsemission spectra before and after the threshold. After thethreshold, a narrow (0.4 nm full width at half maximum)
peak at 413.5 nm was clearly observed. We previously
reported a pulsed operation of a nitride-based VCSEL withthe 3 nm GaInN 5 QWs, showing a threshold current densityof 17 kA=cm2.25) A higher optical conﬁnement factor withthick QWs is one of the reasons for the low threshold current
density.

05FJ11-3© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FJ11 (2016)

4.

T. Furuta et al.

Conclusions

We have achieved a RT CW operation of a nitride-based
VCSEL using 6 nm GaInN 5 QWs and a 4λ-cavity with a
40-pair AlInN=GaN DBR. The threshold current density
under CW operation at RT was 7.5 kA=cm2 with a narrow(0.4 nm) emission spectrum of 413.5 nm peak wavelength.
We also showed that the 6 nm GaInN 5 QWs provided a highoptical conﬁnement factor, a suﬃcient overlap between
electron and hole wave functions, and uniform carrier injections under high current densities.
Acknowledgment

This study was supported by the MEXT Program for theStrategic Research Foundation at Private Universities, 2012–2016.

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Microstructure Analysis of AlGaN on AlN Underlying Layers with Different Threading
Dislocation Densities

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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JE22
http://dx.doi.org/10.7567/JJAP.52.08JE22Microstructure Analysis of AlGaN on AlN Underlying Layerswith Different Threading Dislocation Densities
Kimiyasu Ide1 , Yuko Matsubara1 , Motoaki Iwaya1 , Satoshi Kamiyama1 ,Tetsuya Takeuchi1 , Isamu Akasaki1;3 , and Hiroshi Amano2;3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received October 13, 2012; accepted May 1, 2013; published online August 20, 2013
Using the epitaxial lateral growth technique, we compared the crystallinity and relaxation ratio of 3-m- and 200-nm-thick Al0:5 Ga0:5 N on an AlN
template and AlN grown by epitaxial lateral overgrowth (ELO-AlN), both of which were grown on a sapphire substrate. Although the relaxation
ratios of 3-m-thick Al0:5 Ga0:5 N were almost the same, the misfit dislocation density at the interface and the density of threading dislocations
reaching the surface of Al0:5 Ga0:5 N were significantly different. Also, the increase in the density of newly generated misfit dislocations was found tobe highly dependent on the quality of the AlN underlying layer. We also discuss the difference in the initial growth mode of each Al0:5 Ga0:5 N
sample. # 2013 The Japan Society of Applied Physics

1. Introduction(a)High-crystalline-quality AlGaN is essential for realizing
high-performance UV/deep UV light emitters.1–3) The direct
growth of an AlGaN ﬁlm on a sapphire substrate leads topoor material quality.4–6) When AlGaN is grown on GaN,
because of the substantial lattice mismatch between AlGaNand GaN, tensile stress in AlGaN is released through crack
generation when the critical thickness is exceeded.7) Moreover, the light extraction eﬃciency of UV light-emitting
diodes (UV-LEDs) with an emission wavelength of lower
than 360 nm on GaN is much lower than that of UV-LEDs
on AlN, because GaN absorbs UV light. Thus, the majority
of UV-LEDs8–10) and laser diodes11) have been fabricated ona thick AlGaN/AlN template.
However, there is a lattice mismatch as large as 2.4%
between GaN and AlN. Therefore, the generation of misﬁt
dislocations cannot be avoided when AlGaN is grown on
AlN. It has been reported that the AlN molar fraction densityof dislocations generated in AlGaN on underlying AlN
depends on the AlN molar fraction.12,13) These reportssuggest that the highest dislocation density occurs in
AlGaN with an AlN molar fraction of approximately 0.5.
Furthermore, the microstructure and relaxation process of
Al0:5 Ga0:5 N on AlN with threading dislocation densities(TDDs) of approximately 1  109 and 3  1010 cm2 havebeen reported.14–16) These reports show that the relaxation
mechanism of AlGaN changes with the TDD of the AlN
underlying layer. In the case of Al0:5 Ga0:5 N on AlN witha TDD of approximately 3  1010 cm2 , AlGaN ﬁlms arerelaxed by the inclination of a-type pure edge threading
dislocations (TDs) with Burgers vector b ¼ ð1=3Þh1120i.
AlN ﬁlms with a low dislocation density of approximately 2  108 cm2 grown by epitaxial lateral overgrowth(ELO) have been reported.17) Using this technique, an
improvement in the internal quantum eﬃciency of UVLEDs has been reported.18) As can be seen from the above,
the relaxation process in AlGaN on an AlN template varies
with diﬀerences in the underlying layer such as thediﬀerence in dislocation density. Therefore, there may be
a diﬀerent relaxation process for AlGaN grown on an
ELO-AlN. However, there have been no reports on the(b)Fig. 1. (Color online) Schematic views of (a) 200-nm- and 3-m-thick
Al0:5 Ga0:5 N grown on AlN template, and (b) 200-nm- and 3-m-thick
Al0:5 Ga0:5 N grown on ELO-AlN.

relaxation process of AlGaN on an ELO-AlN underlying
layer.
In this study, we characterized Al0:5 Ga0:5 N on an ELOAlN underlying layer by transmission electron microscopy(TEM), X-ray diﬀraction reciprocal lattice space mapping(XRD mapping) method, and atomic force microscopy(AFM). We compared the crystallinity and relaxation ratioof 3-m- and 200-nm-thick Al0:5 Ga0:5 N on an AlN templateand ELO-AlN, both of which were grown on a sapphire
substrate.
2. Experimental Procedure

Figures 1(a) and 1(b) show schematic views of the sample
structures. All samples were grown on a c-plane sapphire
substrate using a customized high-temperature-growth
metalorganic vapor phase epitaxy (HT-MOVPE) system,
which had a horizontal ﬂow channel conﬁguration and aface-down substrate setup. First, 4-m-thick AlN was grownon the sapphire substrate at a high growth temperature of
1,400  C19,20) at 100 Torr. The total TDD of the AlN ﬁlm
used in this work was approximately 1  109 cm2 . Periodic
grooves along the h1100i
axis, where the width, spacing, and

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Relaxation ratio [%]Jpn. J. Appl. Phys. 52 (2013) 08JE22

K. Ide et al.

AlN template
ELO-AlN

Fig. 3. Plan-view TEM images and TDs at the surface: (a) 3-m-thick
Al0:5 Ga0:5 N grown on AlN template and (b) 3-m-thick Al0:5 Ga0:5 N grownon ELO-AlN.

AlGaN thickness [µm]Fig. 2. (Color online) Relaxation ratio in Al0:5 Ga0:5 N as a function of
Al0:5 Ga0:5 N thickness grown on AlN template and ELO-AlN.

depth are 3, 3, and 1 m, respectively, were formed on
the AlN. Then, the AlN grooves were embedded in an
approximately 20-m-thick ELO-AlN layer. We succeededin growing a low-dislocation-density AlN template with analmost uniform dislocation density of 2  108 cm2 , which
was conﬁrmed by plan-view TEM and from the full widths
at half maximums (FWHMs) of the measured X-ray rocking
curve (XRC). Then, 200-nm- and 3-m-thick Al0:5 Ga0:5 N
layers were grown on the ELO-AlN underlying layer. We
also fabricated 200-nm- and 3-m-thick Al0:5 Ga0:5 N on AlN
templates with an average TDD of 1  109 cm2 as shown in
Fig. 1(a). The AlN molar fraction in AlGaN was estimatedfrom the lattice constants a and c determined by X-ray
diﬀraction measurements.21–23) Cross-sectional TEM specimens of Al0:5 Ga0:5 N samples were prepared using the
focused ion beam technique with low-energy Arþ -ionmilling. Plan-view TEM specimens were prepared fromthe Al0:5 Ga0:5 N samples by standard mechanical polishingand Arþ -ion milling techniques. The microstructure of theAl0:5 Ga0:5 N ﬁlms was analyzed by TEM operating at
300 kV. All observation directions were parallel to theh1100idirection. We also used the g  b method24) in order todetermine the components of the dislocations. The relaxation
ratio and surface condition of each sample were character diﬀraction and byized by XRD mapping around the (2024)
AFM, respectively.

Fig. 4. (Color online) Dark-ﬁeld cross-sectional TEM images of(a) 3-m-thick Al0:5 Ga0:5 N grown on AlN template and (b) 3-m-thick
Al0:5 Ga0:5 N grown on ELO-AlN, taken along the [1100]
projection with

diﬀraction vector of g ¼ ½1120.

3. Results and Discussion

Fig. 5. Plan-view TEM image of 3-m-thick Al0:5 Ga0:5 N grown on ELOAlN near the interface, taken along the [0001] projection.

Figure 2 shows the relaxation ratio in Al0:5 Ga0:5 N as afunction of Al0:5 Ga0:5 N thickness characterized by XRD
 diﬀraction. From this ﬁgure,mapping around the (2024)
there was almost no dependence of the relaxation ratio on
the underlying layer. Moreover, we observed similar lattice
constants a and c for the AlGaN ﬁlms on AlN template and
ELO-AlN. We concluded that AlGaN samples with almost
the same relaxation ratio and AlN molar fraction can beobtained even on diﬀerent underlying layers.
Figure 3 shows bright-ﬁeld plan-view TEM images of(a) 3-m-thick Al0:5 Ga0:5 N on an AlN template and (b)3-m-thick Al0:5 Ga0:5 N on ELO-AlN. From these ﬁgures,
the TDDs of the Al0:5 Ga0:5 N samples were approximately
8:3  109 and 2:6  109 cm2 , respectively. It was clear that
the TDDs at the surface of these two samples were diﬀerent.
That is, although the relaxation ratios of the Al0:5 Ga0:5 N
samples were almost the same, the densities of misﬁt

dislocations of Al0:5 Ga0:5 N on the AlN template and ELOAlN were signiﬁcantly diﬀerent.
Figure 4 shows dark-ﬁeld cross-sectional TEM images of(a) 3-m-thick Al0:5 Ga0:5 N grown on the AlN template and(b) 3-m-thick Al0:5 Ga0:5 N grown on ELO-AlN, taken along
the [1100]
projection with a diﬀraction vector of g ¼ ½1120.
The inclination of a-type dislocations at the AlGaN ﬁlm was
conﬁrmed as shown in Figs. 4(a) and 4(b). Similar results
have been reported in Refs. 14 and 15. Therefore, it wasnecessary to investigate the behavior of misﬁt dislocations at
the interface between Al0:5 Ga0:5 N and AlN. Figure 5 showsa bright-ﬁeld plan-view TEM image at the Al0:5 Ga0:5 N/
ELO-AlN interface extracted from 3-m-thick Al0:5 Ga0:5 N
grown on ELO-AlN. From this ﬁgure, the misﬁt dislocation
density (MDD) generated at the interface between AlGaN

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Jpn. J. Appl. Phys. 52 (2013) 08JE22

K. Ide et al.

Fig. 8. (Color online) AFM images of (a) 200-nm-thick Al0:5 Ga0:5 N
grown on AlN-template and (b) 200-nm-thick Al0:5 Ga0:5 N grown on ELOAlN.
Fig. 6. (Color online) MDD at the interface as a function of thickness of
Al0:5 Ga0:5 N grown on AlN template and ELO-AlN.

Fig. 7. (Color online) TDD at the surface as a function of thickness of
Al0:5 Ga0:5 N grown on AlN template and ELO-AlN.

and AlN in this sample is approximately 2:6  109 cm2 .
In contrast, the MDD of the 3-m-thick Al0:5 Ga0:5 N grownon the AlN template is approximately 7:9  109 cm2 . Thatis, there were signiﬁcant diﬀerences in MDD depending on
the underlying AlN layer.
Figures 6 and 7 show the MDD and TDD of Al0:5 Ga0:5 N
as a function of the Al0:5 Ga0:5 N thickness, respectively.
From these ﬁgures, the TDD of Al0:5 Ga0:5 N on the AlN
template is almost the same as the MDD, but the TDD
of Al0:5 Ga0:5 N on ELO-AlN is lower than the MDD. Next,
we performed AFM observation of the 200-nm-thick
Al0:5 Ga0:5 N on an AlN template and ELO-AlN to investigate
the initial growth modes in Al0:5 Ga0:5 N. Figure 8 shows
AFM images of (a) 200-nm-thick Al0:5 Ga0:5 N grown on an
AlN template and (b) 200-nm-thick Al0:5 Ga0:5 N grown on
ELO-AlN. From these ﬁgures, we conﬁrmed that the growthmodes of AlGaN on the AlN template and ELO-AlN are
signiﬁcantly diﬀerent and involve two-dimensional growth
and three-dimensional growth, respectively. We also characterized the surface structure of ELO-AlN from the AFM
and TEM images. From these images, small steps of
approximately 40 nm at intervals of 6 m were conﬁrmed.
The diﬀerence in growth mode was due to the diﬀerence insurface structure.
We propose a mechanism as illustrated in Fig. 9. The
initial growth of Al0:5 Ga0:5 N on an AlN template is twodimensional, as observed from the initial growth of
Al0:5 Ga0:5 N. In contrast, the initial growth of Al0:5 Ga0:5 N
on ELO-AlN was three-dimensional via the nucleation ofFig. 9. (Color online) Modes of initial growth and growth at a thicknessof 3 m for (a) Al0:5 Ga0:5 N grown on AlN-template and (b) Al0:5 Ga0:5 N
grown on ELO-AlN.

Al0:5 Ga0:5 N. In Al0:5 Ga0:5 N on the AlN template, most of themisﬁt dislocations generated at the Al0:5 Ga0:5 N/AlN interface will thread to the surface since the initial growth modeis two-dimensional. In contrast, lattice relaxation can also
be realized by changing the growth mode to be threedimensional.25,26) Thus, we consider that the three-dimensional growth contributes to the relaxation in Al0:5 Ga0:5 N
on ELO-AlN, in addition to the generation of the misﬁt

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Jpn. J. Appl. Phys. 52 (2013) 08JE22

K. Ide et al.

dislocations. Therefore, the MDD in Al0:5 Ga0:5 N on ELOAlN is less than that of Al0:5 Ga0:5 N on the AlN template.
In addition, the misﬁt dislocations are bent laterally by
three-dimensional growth. Many threading dislocations
do not reach the surface owing to their termination byloop formation. This mechanism explains the signiﬁcantdiﬀerence in the TDD of Al0:5 Ga0:5 N on an AlN templateand ELO-AlN, even though the relaxation ratios of theAl0:5 Ga0:5 N are almost the same.

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p. 165.

4. Conclusions

We compared the crystallinity and relaxation ratio of 3-mand 200-nm-thick Al0:5 Ga0:5 N on AlN templates and ELOAlN grown on sapphire substrates. Although the relaxation
ratios were almost the same, the increases in the misﬁt
dislocation densities of the Al0:5 Ga0:5 N on the AlN templateand ELO-AlN were signiﬁcantly diﬀerent. We also conﬁrmed that the growth modes of Al0:5 Ga0:5 N on the AlN
template and ELO-AlN are signiﬁcantly diﬀerent and exhibit two-dimensional growth and three-dimensional growth,
respectively.

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Acknowledgement

This study was supported by the Ministry of Education,
Culture, Sports, Science and Technology Supported Program
for the Strategic Research Foundation at Private Universities, 2012–2016.

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08JE22-4

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# 2013 The Japan Society of Applied Physics

pssPhys. Status Solidi RRL 7, No. 3, 211– 214 (2013) / DOI 10.1002/pssr.201307023

Analysis of strain relaxation process
www.pss-rapid.com
in GaInN/GaN heterostructure
by in situ X-ray diffraction monitoring during
metalorganic vapor-phase epitaxial growth
Daisuke Iida1, Yasunari Kondo1, Mihoko Sowa1, Toru Sugiyama1, Motoaki Iwaya*, 1, Tetsuya Takeuchi1,
Satoshi Kamiyama1, and Isamu Akasaki1, 2
12

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

Received 18 January 2013, revised 24 January 2013, accepted 25 January 2013
Published online 30 January 2013
Keywords GaInN, heterostructures, in situ XRD, relaxation, pits, misfit dislocation*

Corresponding author: e-mail iwaya@meijo-u.ac.jp

Strain relaxation in a GaInN/GaN heterostructure is analyzed
by combining in situ X-ray diffraction (XRD) monitoring andex situ observations. Two different characteristic thicknessesof GaInN films are defined by the evolution of in situ XRDfrom the full width at half-maximum of symmetric (0002)
diffraction as a function of GaInN thickness. This in situ

XRD measurement enables to clearly observe the critical
thicknesses corresponding to strain relaxation in the
GaInN/GaN heterostructure caused by the formation of surface pits with bent threading dislocations and the generationof misfit dislocations on GaInN during growth.

© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction GaInN is widely used as an active
layer in high-brightness blue and green light-emitting diodes (LEDs) and violet laser diodes [1]. The application of
GaInN films in high-efficiency solar cells is also expected[2]. The bandgap of GaInN alloys ranges from 0.65 to
3.43 eV, making them suitable for the design of highconversion-efficiency multijunction solar cells [3]. So far,
we have succeeded in fabricating GaInN-based solar cells[4–6]. As a result, we have obtained solar cells with up to
4% conversion efficiency by irradiating concentrated
sunlight with intensity up to 300 suns [6]. To realize highperformance GaInN-based solar cells with a high opencircuit voltage and fill factor, it is essential to realize a low
surface pit density of less than 107 cm–2 [4, 5]. Moreover,
GaInN-based green/yellow/red LEDs still have low internal quantum efficiency compared with blue LEDs [7]. Theinternal quantum efficiency of these LEDs decreases withincreasing InN molar fraction in GaInN, which is causedby the generation of misfit dislocations (MDs) at the
GaInN/GaN interface and built-in electrostatic fields [8, 9].
GaInN films with a high InN molar fraction on GaN havealready been used to analyze the relaxation mechanism bywhich defects are formed by strain relaxation [10–14].
However, the sequence of relaxation in the GaInN/GaN
heterostructure remains to be clarified. Moreover, an understanding of the critical layer thickness at which MDs
and surface pits are introduced in the GaInN/GaN heterostructure is essential for the realization of highperformance devices.
In this study, we examined the application of the in situ
XRD observation to GaInN/GaN heterostructures. As a result, we found that this method is useful for analyzing therelaxation process in GaInN/GaN heterostructures by insitu XRD monitoring of the full width at half-maximum(FWHM).
2 Experimental procedure All the samples weregrown by metalorganic vapor phase epitaxy in a face-down
2″ × 3 wafer horizontal-flow reactor. After the growth of a
3-μm-thick GaN template at 1050 °C on c-plane sapphire
covered with a low-temperature buffer layer [15], it wascooled to 755–790 °C. Then, GaInN films of approxi© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysicarrl

212

D. Iida et al.: Analysis of strain relaxation process in GaInN/GaN heterostructure

mately 500 nm thickness with various InN molar fractions
were grown on the GaN template. The threading dislocation (TD) density of the GaN template was approximately
3 × 108 cm–2, which was characterized by plan-view transmission electron microscopy (TEM). We evaluated GaInN
films with symmetric (0002) Bragg diffraction using anin situ XRD system. In situ X-ray was focused on the sample surface using a Johansson curved crystal mirror, andthe diffracted X-ray was detected by a one-dimensionalcharge-coupled device [16, 17]. This system realized theequivalent of a (0002) 2θ/ω scan without any analyzer
crystal in 1 s during rotation of the wafers. In this system,
the tilt component and the distribution of the lattice
constant c can be simultaneously characterized from the
FWHM [18]. We observed the GaInN surface structure anddefects close to the GaInN/GaN heterointerface by scanning electron microscopy (SEM) and TEM, respectively.
Strain relaxation was also evaluated by a typical ex situ
XRD reciprocal space mapping.

X-ray intensity (counts)

3 Results and discussion The in situ XRD spectra of Ga0.87In0.13N/GaN heterostructures are shown in
Fig. 1(a); the InN molar fraction in GaInN can be determined from each peak of the spectrum, and its value corresponds exactly to the value estimated by ex situ XRD. Also,

GaN

105
104

35 nm
80 nm GaInN
170 nm
250 nm103
32.5 33.0 33.5 34.0 34.5 35.02θ (degree)(a)hc1 hc2㻌㻌
1.0

FWHM, ω (degree)dω / dh

0.02

-0.02

hc2

-0.04

0.5

0.0

0.00 saturation

-0.06

0

50 100 150

GaInN thickness, h (nm)

FWHM of underling GaN layer0

100 200 300 400 500

GaInN thickness, h (nm)(b)

Figure 1 (online colour at: www.pss-rapid.com) In situ XRD (a)spectrum and FWHM (b) as a function of Ga0.87In0.13N film
thickness. The experimental FWHMs include the tilt distributionof (0002) diffraction. This indicates that the critical thickness forthe strain relaxation process, hc1, can be obtained by comparing
the results of experiments and simulations. The inset of (b) shows
the derivative of the FWHM with respect to Ga0.87In0.13N film
thickness.
© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

we confirmed that the peak positions of GaInN and GaN
slightly shifted by the influence of the warpage of the sample. Figure 1(b) shows typical in situ XRD FWHMs of theGa0.87In0.13N films during growth. The FWHM of theGa0.87In0.13N films decreases with increasing film thickness,
as determined using a pseudo-Voigt function [16]. The inset in Fig. 1(b) shows the derivative of the FWHM with respect to GaInN thickness. We confirm the existence of adifference between the experimental results and simulation
values. The simulated XRD spectra were obtained using
X’pert Epitaxy 4.0a, a typical X-ray analysis software [19].
The simulated FWHM was then determined using the same
pseudo-Voigt function. Normally, the FWHM decreased tothe same extent as the underlying GaN layer with increasing film thickness; however, this markedly changes for asmall film thickness. The experimental and simulated
FWHMs were shown to similarly decrease with increasing
Ga0.87In0.13N film thickness up to approximately 20 nm.
However, the FWHMs of the Ga0.87In0.13N films withthicknesses exceeding 20 nm tended to deviate from thesimulation values. Then, the FWHM started to saturate at a
GaInN thickness of approximately 100 nm in spite of decreasing simulation values. These behaviors suggest that
the FWHM increases via the generation of MDs and pits
owing to strain relaxation in Ga0.87In0.13N films. Therefore,
we defined hc1 and hc2 to be approximately 20 ± 3 nm and
100 ± 5 nm, respectively.
Figure 2 shows a summary of the growth models alongwith bright-field TEM images, plan-view SEM images,
and XRD (2024) reciprocal space maps of 10-, 35-, 80-,and 110-nm-thick Ga0.87In0.13N films grown on GaN layers.
The TDs almost reached the surfaces of the Ga0.87In0.13N
films grown on the GaN layers in stage (I). We found that
the Ga0.87In0.13N films formed surface pits at the ends of
the TDs with increasing Ga0.87In0.13N film thickness instage (II). In general, surface pits are formed along an inclined {1 101} plane [20]. Here, TDs appear to laterally extend with increasing Ga0.87In0.13N film thickness, as shownin Fig. 2(a) and (b). This suggests that TDs are bent during growth by in-plane compressive stress between
Ga0.87In0.13N and GaN. Then, the surface pits expand withincreasing Ga0.87In0.13N film thickness, as shown in
Fig. 2(f) and (g). It is considered that this promotes thefluctuation of the InN molar fraction in GaInN films by thecomposition pulling effect [11, 12]. Therefore, hc1 is associated with bent TDs, which form surface pits at the endsof TDs. It was clarified that the deviation between the experimental results and simulation values for the FWHM
above a thickness of hc1 increases with increasing apparent
FWHM. Moreover, the density of surface pits increased
with increasing Ga0.87In0.13N film thickness, as shown in
Fig. 2(h). The formation of additional surface pits on the
Ga0.87In0.13N films was mainly caused by TDs induced via
the generation of MDs with strain relaxation in the
Ga0.87In0.13N/GaN heterostructure in stage (III). We foundthat a network of MDs was generated along the ·1 100Ò direction close to the Ga0.87In0.13N/GaN heterointerface; such
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RapidResearch Letter
Phys. Status Solidi RRL 7, No. 3 (2013)

213

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100㻌

dislocations were induced via slips along the {1 123} direction in the {1 122} plane system [21], as shown in Fig. 2(d).
This result suggests that (a + c)-type MDs were induced at
the Ga0.87In0.13N/GaN interface with strain relaxation in
Ga0.87In0.13N. Therefore, although FWHMs by this in situ
XRD do not give any direct information on relaxation,
hc1 and hc2 were indicated as two different critical thicknesses at which the strain relaxation results in the formation of surface pits with bending TDs and the introductionof (a + c)-type MDs, respectively. The GaInN films arealmost perfectly coherently grown on GaN, as indicatedby the results of XRD mapping shown in Fig. 2(i)–(k).
Although strain relaxation was evaluated by conventional
ex situ XRD reciprocal space mapping [21], this method
cannot be used to clearly show the relaxation of bent
TDs with the formation of surface pits on Ga0.87In0.13N
films. At thicknesses exceeding hc2, we only observed therelaxation of MD generation at the heterointerface by exsitu XRD reciprocal space mapping as shown in Fig. 2(l).
That is, relaxation in GaInN can only be confirmed
by XRD reciprocal space maps above a thickness of
hc2. Therefore, in situ XRD monitoring was shown to be auseful method for evaluating hc1 in the thin-film growthstage.

GaInN thickness (nm)Figure 2 (online colour at: www.pss-rapid.com) Summary of several ex situ measurements of Ga0.87In0.13N films with various thicknesses showing plan-view (a)–(d) bright-field TEM and (e)–(h) SEM images, and XRD reciprocal space maps (i)–(l) around (2024)
diffraction. Schematic views of the growth models in stages (I)–(III) are also shown. These stages involve completely coherent growth,
the bending of TDs and the formation of surface pits, and the introduction of MDs, respectively.

10

1
0.0

0.1

0.2

exp: no pits & MDs
exp: pits & bent TDs
exp: MD generation
exp: h c1
exp: h c2
theory: ref. 22
theory: ref. 23
theory: a-type MDs, ref. 24
theory: a+c-type MDs,ref. 24
theory: ref. 25

0.3

InN molar fraction
Figure 3 (online colour at: www.pss-rapid.com) Summary of
GaInN critical thicknesses hc1 and hc2 for various InN molar fractions. The black circles and red triangles indicate the value of
hc1 and hc2, respectively, obtained by in situ XRD measurements.
Blue triangles, circles, and stars indicate ex situ measurement
results for the cases of no pits or misfit dislocations, the appearance of pits and TD bending, and the appearance of misfit dislocations, respectively. The results obtained in situ and ex situ werein good agreement. Theoretical values were taken from Refs.
[22–25].

© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysicarrl

214

D. Iida et al.: Analysis of strain relaxation process in GaInN/GaN heterostructure

We also investigated the values of hc1 and hc2 for various InN molar fractions from the behavior of the FWHMs
by in situ XRD and ex situ measurements, as shown inFig. 3. The results obtained in situ and ex situ were in good
agreement. We also compared our experimental resultswith the results of theoretical calculations in Refs. [22–25].
Theoretical models of the critical thickness are generally
based on lattice relaxation due to the introduction of misfit
dislocations at the heterointerface. However, the experimental results indicate that the MDs are introduced bystrain relaxation after the formation of bent TDs and that
surface pits start to open up at a thickness exceeding hc1.
Therefore, our experimental results are not in good agreement with previous theoretical models. It is considered thatthe difference between the experimental and theoretical results for the generation of MDs in GaInN films is causedby the formation of bent TDs and the opening up of surface
pits. We also observed the behavior of defects in
GaInN/GaN heterostructures with several GaInN thicknesses by TEM. We found that the estimated critical thicknesses were in good agreement with those obtainedby TEM observation. This indicates that the results of insitu and ex situ observations are in relatively good agreement.
4 Summary We observed strain relaxation in GaInN
films during growth by in situ XRD monitoring. The evolutions of XRD peaks and FWHMs were obtained as functions of the GaInN film thickness. Two different featuresof strain relaxation were found: the formation of surface
pits with bending TDs and the introduction of (a + c)-type
MDs. The in situ XRD monitoring system proposed here
enabled us to indicate the lattice properties and observe
several strain relaxation features of GaInN films during
growth. This is considered to be very helpful for growing
high-performance device structures while suppressing theformation of surface pits and MDs.
Acknowledgements This study was supported by the
Ministry of Education, Culture, Sports, Science and Technologyof Japan through the Supported Program for the Strategic Research Foundation at Private Universities, 2012–2016.

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Room-temperature continuous-wave operation of GaN-based vertical-cavity surface-emitting
lasers with n-type conducting AlInN/GaN distributed Bragg reflectors

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Applied Physics Express 9, 102101 (2016)
http://doi.org/10.7567/APEX.9.102101

Room-temperature continuous-wave operation of GaN-based vertical-cavity
surface-emitting lasers with n-type conducting AlInN/GaN distributed
Bragg reﬂectors
Kazuki Ikeyama1, Yugo Kozuka1, Kenjo Matsui1, Shotaro Yoshida1, Takanobu Akagi1, Yasuto Akatsuka1,
Norikatsu Koide1, Tetsuya Takeuchi1*, Satoshi Kamiyama1, Motoaki Iwaya1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering and Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan2*E-mail: take@meijo-u.ac.jp
Received July 18, 2016; accepted August 13, 2016; published online September 1, 2016
The room-temperature continuous-wave operation of a 1.5λ-cavity GaN-based vertical-cavity surface-emitting laser with an n-type conducting
AlInN/GaN distributed Bragg reﬂector (DBR) was achieved. A peak reﬂectivity of over 99.9% was obtained in the n-type conducting AlInN/GaN
DBR so that the current was injected through the DBR for the operation. The threshold current was 2.6 mA, corresponding to the threshold current
density of 5.2 kA/cm2, and the operating voltage was 4.7 V. A lasing spectrum with a peak wavelength of 405.1 nm and a full-width at half maximumof 0.08 nm was also observed. © 2016 The Japan Society of Applied Physics

aN-based vertical-cavity surface-emitting lasers(VCSELs) are expected to be adopted in various
applications, such as retinal scanning displays,
adaptive headlights, and high-speed visible-light communication systems. Recently, the lasing operations of GaN-based
VCSELs with a combination of two dielectric distributed
Bragg reﬂectors (DBRs)1–10) or a combination of a top
dielectric DBR and a bottom undoped GaN-based DBR11–16)
have been demonstrated. Although most of the commercialized infrared VCSELs utilize conducting GaAs-based
DBRs,17) most of the GaN-based VCSELs reported so far
have no other choice but to rely on double intracavity contact
structures due to such insulating DBRs. The double intracavity contacts typically lead to current crowing at current
aperture edges, low optical conﬁnement factors due to a large
cavity length, and complexities of fabrication processes,
resulting in high threshold current densities and low slope
eﬃciencies. If at least an n-type conducting GaN-based DBR
is successfully developed, such as the n-type conducting
GaAs-based DBR, a very short (∼1λ) cavity structure with asimple back-side n-electrode is possible, leading to highperformance and low-cost GaN-based VCSELs.
There are few reports on the n-type conducting GaN-based
DBRs. Arita et al. fabricated a 26-pair n-type Al0.4Ga0.6N=
GaN DBR, showing a peak reﬂectivity of 91% at 400 nm anda series resistance of 180 Ω in a microcavity LED conﬁguration.18) Then, a 20.5-pair n-type AlN=GaN DBR grown on
a SiC substrate by MBE showed a peak reﬂectivity of 99% at
450 nm and a speciﬁc series resistance of 2 × 10−3 Ω cm2.19)
Also, a 40-pair n-type Al0.12Ga0.88N=GaN DBR showed apeak reﬂectivity of 91.6% at 368 nm and a bulk resistivityof 0.52 Ω cm.20) However, so far, there has been no reporton the laser operation of a GaN-based VCSEL with a conducting DBR. We have been developing insulating andn-type conducting Al0.82In0.18N=GaN DBRs.21,22) AlInN is
lattice-matched with GaN, so that the AlInN=GaN DBRs
have potentials of higher material quality and reﬂectivity thanthe AlGaN=GaN DBRs,23,24) possibly comparable to the
GaAs-based DBRs.25) A negative aspect of AlInN was thevery low growth rate, but we have optimized the growth conditions for AlInN with a growth rate of over 0.5 µm=h.21)
Regarding the conductivity, it must be considered that thelarge polarization charges induced at the AlInN=GaN inter-

G

(a)

(b)

Fig. 1. (a) Schematic of a 10-pair Si-doped AlInN=GaN DBR structure forvertical current injection and (b) a Si-doping proﬁle in a pair of AlInN=GaN
layers.

faces markedly aﬀect the vertical conductivity through theinterfaces.26) We then proposed and developed a new conceptof modulation doping for a nitride-based DBR to neutralize
the large polarization charges with ionized donors.22,27)
In this study, we developed n-type conducting AlInN=GaN
DBRs and demonstrated the room-temperature continuouswave (CW) operation of a GaN-based VCSEL utilizing then-type conducting AlInN=GaN DBR.
We prepared four diﬀerent samples on free-standing (0001)
n-GaN substrates (carrier concentration: 1 × 1018 cm−3, thickness: 330 µm) by metalorganic vapor phase epitaxy. The
ﬁrst three samples were AlInN=GaN DBRs. As shown in
Fig. 1(a), a 10-pair Si-doped AlInN=GaN DBR with a 5 nm
n+-GaN contact layer was prepared in order to investigate
vertical current injection. A 1.0 µm n-GaN buﬀer layer witha Si concentration of 6.0 × 1018 cm−3 was grown before the
DBR growth. The thicknesses of AlInN and GaN were 40–45
nm, designed as a peak reﬂectivity wavelength of 405 nm.
After the DBR growth, the 5 nm n+-GaN contact layer with a
Si concentration of 1 × 1020 cm−3 was capped. On the other
hand, as described in Fig. 1(b), a speciﬁc modulation dopingwith Si was formed in the DBR based on our previous
results.22) Although the Si concentrations in most regions
were in the range of 1018–1019 cm−3, a very high Si concentration (6 × 1019 cm−3) was doped in 15 nm regions including

102101-1© 2016 The Japan Society of Applied Physics

Appl. Phys. Express 9, 102101 (2016)

K. Ikeyama et al.

Fig. 3. J–V characteristics of a 10-pair Si-doped AlInN=GaN DBR.
Fig. 2. Schematic of a 1.5λ-cavity GaN-based VCSEL with verticalcurrent injection through the n-type conducting AlInN=GaN DBR.

the top AlInN interfaces with GaN at which a large negative
polarization charge concentration (3 × 1013 cm−2) was generated. Note that Si could be ionized and positively charged
with the concentration of 6 × 1019 cm−3 in the 15 nm regions
corresponding to a sheet concentration of 9 × 1013 cm−2.
Therefore, the Si concentration around the interface should be
suﬃciently high to completely neutralize the high negative
polarization charge concentration. This is a concept that we
previously proposed, suﬃciently eliminating energy spikes
in the band diagram at the GaN=AlInN interfaces due to thepolarization charges in order to obtain a low vertical resistancethrough the interfaces.22,27) A 60-µm-diameter mesa wasformed by dry etching, and Cr=Ni=Au contacts were deposited on the top of the mesa and on the back side of thesubstrate. Current density–voltage (J–V ) characteristics undera vertical current injection through the DBR were measured.
The second DBR sample was a 46-pair Si-doped AlInN=
GaN DBR for reﬂectivity and surface morphology measurements. The same modulation doping with Si was performed,
but no contact layer was grown on the DBR. For comparison,a 40-pair undoped AlInN=GaN DBR, which was used in
our previous VCSELs showing room-temperature CW
operations,15,16) was also prepared as the third DBR sample.
Reﬂectivity spectra of the DBRs were measured with anabsolute reﬂectivity measurement system, and the surface
morphology was observed by atomic force microscopy.
The fourth sample was a GaN-based VCSEL with the46-pair Si-doped AlInN=GaN DBR as shown in Fig. 2. After
growing the AlInN=GaN DBR, a 1.25λ-cavity containing
50 nm n-GaN, ﬁve 3 nm GaInN=6 nm GaN quantum wells, a
20 nm p-Al0.2Ga0.8N electron blocking layer, 60 nm p-GaN,
and a 10 nm p+-GaN contact were continuously grown on the
DBR. Next, a 39 µm mesa of 200 nm height was formed on
the wafer for device isolation by dry etching. Note that the
Si-doped AlInN=GaN DBR was hardly etched out in the
VCSEL structure. Then, 10 nm SiO2 with an 8-µm-diameter
aperture, a 20 nm indium–tin oxide (ITO) p-electrode, and a
32 nm Nb2O5 spacer layer were deposited, forming a 1.5λcavity. The emission area was deﬁned by the 8-µm-diameter
aperture of SiO2. Next, an 8-pair Nb2O5=SiO2 top DBR,
showing a 99.9% reﬂectivity, was also deposited. Finally,a p-pad on the ITO layer and an n-electrode on the n-GaN
substrate back side were deposited. Current–voltage–light
output power (I–V–L) characteristics were measured under

CW operation at room temperature. Emission spectra andpolarization characteristics were also measured. The light
output power was coupled to a ﬁber cable in the abovemeasurements.
Figure 3 shows the J–V characteristics of the 10-pairSi-doped AlInN=GaN DBR for the vertical current injection.
The positive voltage in the ﬁgure means that a positive
voltage was applied to the top n-electrode on the mesa.
A linear relationship between current density and voltage
was clearly observed in both forward and reverse directions.
A speciﬁc series resistance of 1.7 × 10−4 Ω cm2 was obtained.
Considering the number of pairs and the contribution ofn-electrode contact resistances, the resistance of a 46-pair
DBR is estimated to be 7.8 × 10−4 Ω cm2 or lower, which
is lower than that previously reported for the GaN-based
DBR,19) but still higher than that reported for a p-type
GaAs-based DBR (6.2 × 10−5 Ω cm2).17)
Figures 4(a) and 4(b) show wide-range and narrow-range
reﬂectivity spectra of the 46-pair Si-doped and 40-pairundoped AlInN=GaN DBRs, respectively. A large stopband
width (22 nm) above a 90% reﬂectivity and many sharp side
lobes were consistently observed from both the DBRs as
shown in Fig. 4(a). Furthermore, the 46-pair Si-doped AlInN=
GaN DBR showed a peak reﬂectivity of over 99.9% at 404
nm, which was clearly higher than that of the 40-pair undoped
AlInN=GaN DBR. These results suggest that our modulation
doping with Si in the AlInN=GaN DBR does not adversely
aﬀect the reﬂectivity characteristics. The higher reﬂectivity of
the 46-pair DBR was mainly due to 6 additional pairs. We also
observed a well-ordered surface morphology of the Si-doped
DBR covered with atomic layer steps in Fig. 5, suggestingthat a small surface roughness (RMS value: 0.08 nm) led to alow scattering loss as a reﬂector. As indicated above, we haveobtained the 46-pair Si-doped AlInN=GaN DBR showing
the very high reﬂectivity and reasonably good electricalconductivity simultaneously.
Figures 6(a) and 6(b) show the I–L–V characteristics andemission spectra of the 1.5λ-cavity GaN-based VCSEL with
the n-type conducting AlInN=GaN bottom DBR under CW
operation at room temperature. A marked increase in light
output power was clearly observed at 2.6 mA, and also a very
narrow emission spectrum (FWHM: 0.08 nm) at 405.1 nm
was observed. In Fig. 6(b), a redshift of the emission peak
was also observed with an increase of the injected current. At
this moment we believe the shift was caused by heating due
to a low thermal conductivity of AlInN layers.28) Figure 7
shows light output power as a function of polarizer angle at102101-2© 2016 The Japan Society of Applied Physics

Appl. Phys. Express 9, 102101 (2016)

K. Ikeyama et al.

(a)
(a)

(b)
Fig. 6. (a) I–L–V characteristics and (b) emission spectra of the 1.5λcavity GaN-based VCSEL with the 46-pair n-type conducting AlInN=GaN
DBR.

(b)
Fig. 4. (a) Wide-range and (b) narrow-range reﬂectivity spectra of a 46pair n-type conducting AlInN=GaN DBR and a 40-pair undoped AlInN=GaN
DBR.

Fig. 7. Light output power of a VCSEL as a function of polarizer angle at
5.0 mA.
Fig. 5. 2 × 2 µm2 AFM image of a 46-pair n-type conducting AlInN=GaN
DBR.

5.0 mA, indicating that the emission light was linearlypolarized. Thus, we concluded that the GaN-based VCSEL
with the vertical current injection through the conducting
bottom DBR showed a room-temperature CW operation witha threshold current of 2.6 mA, corresponding to a thresholdcurrent density of 5.2 kA=cm2. The operating voltage anddiﬀerential device resistance at the threshold current were
4.7 V and 250 Ω, respectively. Previously, we demonstrated
the 4λ-cavity GaN-based VCSEL containing the double
intracavity contacts with the insulating AlInN=GaN DBR.
The VCSEL showed a 7.5 kA=cm2 threshold current density
and a 4.8 V operating voltage.16) Thus, we concluded that thecharacteristics of our GaN-based VCSEL with the n-typeconducting bottom DBR were comparable to or even better
than those of the GaN-based VCSEL with the insulating
DBR. Further improvements of the characteristics of theGaN-based VCSELs with the n-type conducting DBRs are
expected with the further optimization of the modulation
doping of Si in terms of not only resistivity but also theabsorption and appropriate design=control of layer thicknesses in the VCSELs.
In summary, a room-temperature CW operation of the
1.5λ-cavity GaN-based VCSEL with an n-type conducting
AlInN=GaN bottom DBR was achieved. The n-type conducting AlInN=GaN DBR simultaneously showed an ohmic
behavior and a high reﬂectivity (99.9%). As a result, theGaN-based VCSEL showed a threshold current density of102101-3© 2016 The Japan Society of Applied Physics

Appl. Phys. Express 9, 102101 (2016)

K. Ikeyama et al.

5.2 kA=cm2 and an operating voltage of 4.7 V at a lasing
wavelength of 405.1 nm. These results indicate the feasibilityof high-performance and low-cost GaN-based VCSELs withvertical current injections through conducting DBRs.
Acknowledgment This study was supported by the MEXT-Supported
Program for Strategic Research Foundation at Private Universities, 2012–2016.

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Control of the Detection Wavelength in AlGaN/GaN-Based Hetero-Field-Effect-Transistor
Photosensors

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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JF02
http://dx.doi.org/10.7567/JJAP.52.08JF02

Control of the Detection Wavelength in AlGaN/GaN-Based
Hetero-Field-Effect-Transistor Photosensors
Mami Ishiguro1 , Kazuya Ikeda1 , Masataka Mizuno1 , Motoaki Iwaya1 , Tetsuya Takeuchi1 ,Satoshi Kamiyama1 , and Isamu Akasaki1;2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
E-mail: 123434004@ccalumni.meijo-u.ac.jp
2

Received October 12, 2012; accepted February 13, 2013; published online May 31, 2013
We examined the control of the detection wavelength in AlGaN/GaN-based hetero-field-effect-transistor (HFET) photosensors. The detection
wavelength of these devices can be controlled by using the p-GaInN optical gate or inserting a GaInN channel layer between AlGaN and GaN. In
addition, the photosensitivity of AlGaN/GaN HFET photosensors with a p-GaInN optical gate was more than two orders of magnitude higher thanthat of the AlGaN/GaN HFET photosensor with a GaInN channel layer. Moreover, the photosensitivity of the AlGaN/GaN HFET photosensor witha p-GaInN optical gate greatly surpassed those of commercially available Si pin and Si avalanche photodiodes, and was comparable to those ofphotomultiplier tubes. # 2013 The Japan Society of Applied Physics

1. Introduction

Since group-III-nitride semiconductors, i.e., InN, GaN, AlN,
and their alloys, have direct wide band gaps ranging from0.67 to 6.2 eV,1–3) they are suitable materials for opticaldevices for light in the UV to near-infrared range, which
includes the entire visible region. High-performance emitters, such as UV-,4,5) blue-, green-, and white-light-emitting
diodes6) and UV-7) violet-,8) blue-,9) and green10)-laser
diodes, have already been achieved using these materials.
In contrast, the physical properties of nitride semiconductors are expected to be applied to other devices, such ashetero-ﬁeld-eﬀect-transistors11,12) and solar cells.13,14) Moreover, the nitride-based photosensors in the UV-region havealso attracted attention for such applications as ﬂame
detection and biosensors.14–19) The nitride-based photosensors are also promising as a technology for frontier
applications, such as wireless visible-light communication.
Wireless visible-light communication has attracted attention,for example, for use in the optical bus of computers and for
wireless access in vehicles, because it has major advantages
such as no eﬀect of electromagnetic radiation, low cost,and high directionality. In wireless visible-light communication, Si-based photosensors, such as pin diodes, areused. However, the communication speed of such a system
is limited to several hundred kbps because of the low
performance of the photosensors. To increase the communication speed, the photosensor used must have a highphotosensitivity, a high ratio of photocurrent to dark current(high S/N ratio), and a high response speed.
Recently, we have reported novel high-performance
nitride-based AlGaN/GaN hetero-ﬁeld-eﬀect-transistor-type(HFET-type) photosensors with a p-GaN optical gate.20)
These photosensors employ a two-dimensional electron
gas at the heterointerface between AlGaN and GaN as ahighly conductive channel with a high electron mobility.
The photosensitivity of these devices greatly surpassed thoseof commercially available Si pin and Si avalanche photodiodes, and was comparable to those of photomultipliertubes.21)
Figure 1 shows schematic diagrams of the operating
principle of this device. Such devices take advantage of the

Depletion layer
p-GaN

Light
p-GaN

2DEG

u-AlGaN

u-AlGaNu-GaN
LT-Buffer Layeru-GaNLT-Buffer Layerc-plane Sapphirec-plane Sapphire

(a)

(b)

Fig. 1. (Color online) Schematic diagrams of the operating principle of
HFET-type photosensor with p-GaN optical gate (a) with no irradiation oflight and (b) with irradiation of light.

two-dimensional electron gas (2DEG) by means of the
AlGaN/GaN heterostructure and depletion layer formed by
the p-GaN optical gate. In Fig. 1(a), a 2DEG is formed at the
AlGaN/GaN heterointerface region without the p-GaN.
Thus, contact and sheet resistances in this region are verysmall. In contrast, because the depletion layer is formed
immediately beneath the p-GaN, sheet resistance in thisregion is very high. According to the concept of the voltage
divider circuit, most of the voltage is applied immediately
beneath the p-GaN. Because this distance is several
micrometers, an electric ﬁeld as high as 104 V/cm is applied
only immediately beneath the p-GaN, when a voltage ofseveral volts is applied between the anode and cathodeelectrodes. In addition, the electron mobility of the 2DEG
in the AlGaN/GaN heterointerface is approximately 1000
cm2 V 1 s 1 . When electron–hole pairs are formed in the
p-GaN optical gate by irradiating light, the conduction bandedge in the AlGaN/GaN heterostructure is considered tochange. The drift current density is determined as the
product of the elementary charge, electron concentration,
electron mobility, and electric ﬁeld. Accordingly, a large
photocurrent can be obtained by irradiating weak light to the
AlGaN/GaN heterostructure, as shown in Fig. 1(b). Therefore, this structure can be expected to have an extremely

08JF02-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JF02

M. Ishiguro et al.
8µm

high photosensitivity and a high S/N ratio. However, there
are many unclear points in the operating principle of thesedevices. In particular, control of the detection wavelength inthese devices is unknown.
In this study, we examined the control of the detection
wavelength in the AlGaN/GaN-based HFET photosensors.
We also discuss the performance of these devices.

Anode(Ti/Al/Ti/Au)

2µm

Cathode(Ti/Al/Ti/Au)

u-Al0.2Ga0.8N:12nm
p-Ga0.86In0.14N(110nm)u-GaN:~3µm
LT-Buffer Layer

2. Experimental Procedure

Figure 2 shows a schematic view of the HFET-type
photosensor with a p-layer optical gate. The device wasgrown by metalorganic vapor phase epitaxy on a sapphire(0001) substrate covered with a low-temperature-deposited(LT-) buﬀer layer.22) In this study, we fabricated threesamples for controlling the detection wavelength. In type I[Fig. 2(a)], we fabricated an AlGaN/GaN HFET-type
photosensor with a p-GaInN optical gate. This structure isa stack comprising a 3.0-m-thick unintentionally doped
GaN (u-GaN) layer, a 12-nm-thick unintentionally doped
Al0:2 Ga0:8 N (u-Al0:2 Ga0:8 N) barrier layer, and a 110-nmthick Mg-doped p-Ga0:86 In0:14 N layer with a Mg concentration of approximately 3  1019 cm 3 . In type II [Fig. 2(b)],we fabricated the AlGaN/GaN HFET with a GaInN channel
layer inserted between AlGaN and GaN with a p-GaN
optical gate. This structure is stacked with a 3.0-m-thicku-GaN layer, a 2-nm-thick u-Ga0:90 In0:10 N channel layer, a12-nm-thick u-Al0:2 Ga0:8 N barrier layer, and a 110-nm-thickMg-doped p-GaN layer with a Mg concentration of
approximately 3  1019 cm 3 . In type III [Fig. 2(c)], we
fabricated the AlGaN/GaN HFET with a 120-nm-thick
p-GaN optical gate as a reference. The AlN molar fraction
and thickness of the AlGaN barrier layer were determinedby X-ray diﬀraction 2=! scans with (0002) diﬀraction. The
hole concentrations in the p-Ga0:86 In0:14 N and p-GaN opticalgates were approximately 2  1018 and 1  1018 cm 3 atroom temperature (RT), respectively.
After the activation of Mg acceptors by heating at 525  C
for 4 min in N2 , mesa isolation was performed by Cl2
reactive ion etching (RIE). Next, the p-layers were etched by
RIE, except in the light-receiving area. Then, Ti/Al/Ti/Au
was deposited as the anode and cathode contacts on the
u-Al0:2 Ga0:8 N layer. The length and width of the p-GaInN
optical gate were 2 and 100 m, respectively, and theinterval between the anode and cathode contacts was 8 m.
No antireﬂection coating was used for any of the devices inthis study. The photocurrent at each light wavelength wasmeasured with a semiconductor parameter analyzer (Agilent
Technologies HP-4155C). The photosensitivity of eachphotosensor was calculated from the photocurrent measured
at light wavelength intervals of 5 nm using a spectroscopewith a Xe lamp at RT. The intensity of the spectral radiance
was measured at each wavelength. The monochromatic lightirradiation densities of the spectroscope-measured photosensitivity ranged from 200 to 220 W/cm2 . Since the lightreceiving area was 200 m2 , the amount of received light
was very small and equivalent to approximately several
hundred picowatts.
3. Results and Discussion

Figure 3 shows the monochromatic photosensitivity ofeach HFET-type photosensor with a driving voltage between

c-plane Sapphire
(a)
8µm

Anode
(Ti/Al/Ti/Au)

p-GaN(110nm)Cathode(Ti/Al/Ti/Au)

2µm

u-Al0.2Ga0.8N:12nm
u-Ga0.9In0.1N;2nm
u-GaN:~3µm
LT-Buffer Layerc-plane Sapphire
(b)

Anode
(Ti/Al/Ti/Au)

p-GaN(120nm)8µµm
2µm

Cathode(Ti/Al/Ti/Au)

u-Al0.2Ga0.8N:12nm
u-GaN:~3µm
LT-Buffer Layerc-plane Sapphire(c)Fig. 2. (Color online) Schematic views of device structures: (a) HFETtype photosensor with p-GaInN optical gate, (b) HFET-type photosensor
with GaInN channel layer, and (c) HFET-type photosensor with p-GaN
optical gate.

the anode and cathode (VAC ) of 5 V obtained using a Xelamp and a spectrometer. Figure 4 also shows the normalized monochromatic photosensitivity of each photosensor.
The absorption edge wavelengths of the photosensors withthe p-GaInN optical gate, the GaInN channel, and the p-GaN
optical gate were approximately 450, 410, and 375 nm,
respectively, as shown in Fig. 4. From this result, weconcluded that the detection wavelength of the photosensitivity of the HFET-type photosensor can be controlled bycontrolling the InN molar fraction in the p-GaInN opticalgate or the GaInN channel layer. In addition, the photosensor
with the p-Ga0:8 In0:2 N optical gate exhibits photosensitivity
up to a longer wavelength of 450 nm. This is considered
to be due to the diﬀerence in the InN molar fraction in the
p-GaInN optical gate and the GaInN channel. Moreover,the maximum photosensitivity of the photosensor with the
p-GaInN optical gate is three orders of magnitude higher
than that of the GaInN-channel sample. Since the mobility of
the HFET with a GaInN channel is about two orders of

08JF02-2# 2013 The Japan Society of Applied Physics

Photosensitivity [A/W]

Jpn. J. Appl. Phys. 52 (2013) 08JF02

10

6

10

5

10

4

10

3

10

2

10

1

Table I. Photosensitivity, detection wavelength, and dark current ofcommercial Si pin photodiode, Si APD, photomultiplier tubes, and HFET
photosensors in this work.

p-GaN optical gate
p-GaInN optical gate
GaInN channel

Structure

0

RT

-1

VDS=5V1010

M. Ishiguro et al.

HFET-type photosensor
with p-GaInN optical gate
HFET-type photosensor
with GaInN channel layer
HFET-type photosensor
with p-GaN optical gate
Si p–i–n photodiode19)19)

Si APD

19)

Photomultiplier tube

300 400 500 600
Wavelength [nm]Normalized photosensitivity

Fig. 3. (Color online) Monochromatic photosensitivity of each HFETtype photosensor.

1.0
0.8

p-GaN optical gate
p-GaInN optical gate
GaInN channel

0.6
0.4
0.2
0.0
300 350

VAC = 5 V
RT

400 450500

Detection
wavelength(nm)

Maximum
photosensitivity
(A/W)

Dark
current(A)450

4  105

10 5

410

1  102

10 5

375

4  104

10 9

320–1100

4  10 1

2:0  10 10

320–10001

4  10

5:0  10 8

185–900

5

1:0  10 8

9  10

we believe that the dark current is possible to be reduced by
optimizing the device structure.27,28)
In addition to the high sensitivity, a high-speed response
can be expected, because the response of this device isthought to be determined by the transit time of the 2DEG
under the optical gate or the photocarrier recombination timein the p-GaInN optical gate, which are both estimated to beless than 1 ns.29,30) The dynamic performance of this device
will be examined in the near future.
4. Conclusions

550

Wavelength [nm]Fig. 4. (Color online) Normalized monochromatic photosensitivity ofeach HFET-type photosensor.

magnitude lower than that of the AlGaN/GaN HFET,23) weinferred that such a diﬀerence in photosensitivity indeed
exists in these devices. Therefore, we concluded that the
p-GaInN optical gate is the best structure for controllingthe detection wavelength of the AlGaN/GaN HFET-type
photosensors. The photocurrent-to-dark current (S/N) ratiosof the p-GaInN optical gate and the GaInN channel were on
the order of 102 and 104 , respectively.
Table I summarizes the photosensitivity, detection wavelength, and dark current of each photosensor. We also
included the photosensitivity of our photosensors andthose of commercial Si p–i–n photodiodes (Hamamatsu
Photonics S2387-1010R),24) Si avalanche photodiodes(APD; Hamamatsu Photonics S8664-55),25) and photomultiplier tubes (Hamamatsu Photonics H7421-40, H7421-50,
R928, and R3896).26) The photosensitivity of the HFET-type
photosensor with a p-GaInN optical gate greatly surpassed
those of the Si pin photodiode and Si APD and wascomparable to those of the photomultiplier tubes. In
addition, the dark currents of the HFET-type photosensor
using a p-GaInN optical gate and GaInN channel layer are
higher than those of other devices. We thought that this iscaused by an insuﬃcient depletion layer in the 2DEG in a
p-GaInN optical gate and GaInN channel devices. Therefore,

We fabricated high-performance HFET-type photosensors
for the detection of visible light. The detection wavelengthof the photosensitivity of the HFET-type photosensor canbe controlled by controlling the InN molar fraction in the
p-GaInN optical gate or GaInN channel layer. Since themaximum photosensitivity of a photosensor with the
p-GaInN optical gate is three orders of magnitude higher
than that of the GaInN-channel sample, we concluded thatthe p-GaInN optical gate for controlling the detection
wavelength of the AlGaN/GaN HFET-type photosensors isthe best structure. The photosensitivity of the HFET-type
photosensors with a p-GaInN optical gate greatly surpassed
those of the commercially available Si pin photodiode and
Si APD and was comparable to those of photomultipliertubes.
Acknowledgments

This study was supported by the Ministry of Education,
Culture, Sports, Science and Technology of Japan, throughthe Supported Program for the Strategic Research Foundation at Private Universities, 2012–2016, and the Knowledge
Cluster Initiative (Second Stage) of the Tokai Region
Nanotechnology Manufacturing Cluster.

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08JF02-3# 2013 The Japan Society of Applied Physics

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# 2013 The Japan Society of Applied Physics

pssPhys. Status Solidi RRL 7, No. 3, 215– 217 (2013) / DOI 10.1002/pssr.201206483

Nitride-based hetero-field-effecttransistor-type photosensors
with extremely high photosensitivity

www.pss-rapid.com

Mami Ishiguro1, Kazuya Ikeda1, Masataka Mizuno1, Motoaki Iwaya*, 1, Tetsuya Takeuchi1,
Satoshi Kamiyama1, and Isamu Akasaki1, 2
12

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

Received 13 November 2012, revised 16 January 2013, accepted 17 January 2013
Published online 22 January 2013
Keywords GaInN, photosensors, HFET, field-effect transistors, two-dimensional electron gases*

Corresponding author: e-mail iwaya@meijo-u.ac.jp

AlGaN/GaN hetero-field-effect-transistor-type (HFET-type)photosensors are fabricated with a p-GaInN optical gate forthe detection of visible light. These photosensors employ a
two-dimensional electron gas at the heterointerface between
AlGaN and GaN as a highly conductive channel with a highelectron mobility. By changing the InN molar fraction in the

p-GaInN optical gate, the wavelength range of the photosensitivity of the HFET-type photosensors can be controlled. The
photosensitivity of the AlGaN/GaN HFET-type photosensors
with a p-GaInN optical gate greatly surpassed those of commercially available Si pin and Si avalanche photodiodes, andwas comparable to those of photomultiplier tubes.

© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction Wireless visible-light communication
has attracted attention for use in the optical bus in computers, for wireless access in vehicles, and so forth, because it has major advantages such as no effect of electromagnetic radiation, low cost, and high directionality. For
these applications, the photosensor used must have a highphotosensitivity, a high ratio of photocurrent to dark current (high S/N ratio), and a high response speed. However,no appropriate devices with a high photosensitivity and ahigh S/N ratio in the visible region have been developed.
For instance, although a photomultiplier tube is able to detect light with an externally high photosensitivity, it isbreakable, expensive, has a short lifetime, and requires ahigh operation voltage exceeding several hundred volts.
Moreover, the photosensitivities of Si pin and avalanche
photodiodes are up to four orders of magnitude lower thanthat of a photomultiplier tube.
Since group-III-nitride semiconductors, i.e., InN, GaN,
AlN, and their alloys, have direct wide band gaps ranging
from 0.67 to 6.2 eV [1–3], they are suitable materials for
optical devices ranging from UV to near-infrared light,
which includes the entire visible region. High-performance
emitters, such as UV- [4], blue-, green-, and white-light-

emitting diodes and violet-, blue-, and green-laser diodes[5], have already been achieved using these materials. In
contrast, although there have been several reports of nitridebased photodetectors, such as photoconductors [6], Schottky photovoltaic detectors [7], pin photodiodes [8], solar
cells [9, 10], and photo field-effect transistors [11], no appropriate devices with both a high photosensitivity and ahigh S/N ratio have been realized. In addition, with the exception of solar cells, most devices only exhibit photosensitivity in the UV region, and cannot detect light in the visible
region. Therefore, there are no suitable photosensors forapplication to wireless visible-light communication.
In this study, we fabricated high-performance nitridebased AlGaN/GaN heterostructure-field-effect-transistortype (HFET-type) photosensors with a p-GaInN opticalgate for the detection of visible light. These photosensors
employ a two-dimensional electron gas (2DEG) at the heterointerface between AlGaN and GaN as a highly conductive channel with a high electron mobility. In addition, ap-GaInN layer ensures a high photosensitivity in the visible region and a high S/N ratio. The obtained results arevery promising for the development of high-performance
photosensors for wireless visible-light communication.
© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

216

M. Ishiguro et al.: Nitride-based hetero-field-effect-transistor-type photosensors

diation of weak light in the AlGaN/GaN heterostructure.
Therefore, this structure can be expected to have an extremely high photosensitivity and a high S/N ratio.
The photosensitivity of each photosensor was calculated from the photocurrent measured at light wavelengthintervals of 5 nm using a spectroscope with a Xe lamp at
RT. The photocurrent at each light wavelength was measured by a semiconductor parameter analyzer (Agilent
Technologies, Inc.: HP-4155B). The monochromatic lightirradiation densities of the spectroscope-measured photosensitivity ranged from 200 to 220 μW/cm2. Since thelight-receiving area was 200 μm2, the amount of received
light was very small and equivalent to approximately several hundred picowatts.
3 Results and discussion Figure 2(a) and (b), respectively, show the monochromatic photosensitivity of
HFET-type photosensors with p-Ga0.86In0.14N and p-GaN
optical gates with a driving voltage between the anode andthe cathode (VAC) of 5 V obtained using a Xe lamp andspectrometer. Figure 3 shows the normalized monochromatic photosensitivity of each photosensor. From thesefigures, although there are differences in the photosensitivity of each device, the photosensors can achieve a high external photosensitivity, exceeding 104 A/W. Also, dark current densities of the photosensors with the p-Ga0.86In0.14N
and p-GaN optical gates of approximately 1 μA/mm andlower than 10 nA/mm, respectively, were obtained at
VAC = 5 V. The S/N ratios of the photocurrent to dark current of each sample were on the order of 102 and 104,
respectively. The dark current of such HFET-type photosensors using a p-GaInN optical gate is higher than thoseusing a p-GaN optical gate. We believe that this is causedby an insufficient depletion of the 2DEG in the case of the
p-GaInN optical gate and optimization should allow a further reduction of the dark current. Although the device
structure needs to be optimized to maximize the photosensitivity and S/N ratio of the device, we have shown that aphotosensor with a high external photosensitivity can beachieved. In addition, Fig. 3 shows normalized monochromatic photosensitivities of HFET-type photosensors.
From this figure, the absorption edge wavelengths of thephotosensors with p-Ga0.86In0.14N and p-GaN optical gates

VAC = 5 V
5© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

RT

4x10

5

3x10

5

2x10

5

1x10

0
300

Figure 1 (online colour at: www.pss-rapid.com) Schematic viewand top-view microscopic image of HFET-type photosensor.

4

5x10

5

5x10

Photosensitivity [A/W]

2 Experimental procedure Figure 1 shows a schematic view and a microscopic top-view image of an HFETtype photosensor with p-GaInN optical gate. The device
was grown by metalorganic vapor phase epitaxy on a sapphire (0001) substrate covered with a low-temperaturedeposited (LT-)buffer layer [12]. After depositing the
20-nm-thick LT-buffer layer at 500 °C, a 3.0-μm-thickunintentionally doped GaN (u-GaN) layer, a 12-nm-thickunintentionally doped Al0.2Ga0.8N (u-Al0.2Ga0.8N) barrier
layer, and a Mg-doped p-GaInN layer with a Mg concentration of approximately 3 × 1019 cm–3 were grown. We
fabricated two devices with p-GaInN layer having InN molar fractions of 0.14 and 0. The AlN molar fraction andthickness of the AlGaN barrier layer were determined by
X-ray diffraction 2θ/ω scans with (0002) diffraction. The
hole concentrations in the p-Ga0.86In0.14N and p-GaN optical gate, which were determined by Hall effect measurement, were approximately 2 × 1018 and 3 × 1018 cm–3 atroom temperature (RT), respectively.
After the activation of Mg acceptors by heating to
525 °C for 4 min in N2, mesa isolation was performed by
Cl2 reactive ion etching (RIE). Next, the p-GaInN opticalgate was etched by RIE, except in the light-receiving area.
Then, Ti/Al/Ti/Au was deposited as the anode and cathodecontacts on the u-Al0.2Ga0.8N layer. The length and thewidth of p-GaInN optical gate were 2 μm and 100 μm, respectively, and the interval between the anode and cathodecontacts was 8 μm. No antireflection coating was used forany of the devices in this study.
In these devices, in absence of any p-layer, a 2DEG isformed at the AlGaN/GaN heterointerface region. Thus,
contact and sheet resistances in this region are very small.
The electron mobility is approximately 1000 cm2/Vs. However, in presence of a p-GaInN layer directly on top of thedepletion layer, the sheet resistance in this region isvery high. Therefore, whenever, a voltage of several volts
is applied between the anode and cathode electrodes, anelectric field as high as 104 V/cm acts across the depletion
layer. Next, however, when by the irradiation of light onto
the p-GaInN optical gate, electron–hole pairs are generated
and the conduction band edge in the AlGaN/GaN heterostructure is considered to change. The drift current density depends on the product of the elementary charge, electron concentration, electron mobility, and electric field.
Accordingly, a large photocurrent can be obtained by irra-

350

400

450

500

㻌
Photosensitivity[A/W]solidistatusphysicarrl

VAC = 5 V
RT

4

4x10

4

3x10

4

2x10

4

1x10

0
300

350

400

Wavelength [nm]Wavelength [nm](a) p-GaInN(b) p-GaN

450

Figure 2 (online colour at: www.pss-rapid.com) Monochromatic photosensitivity of HFET-type photosensors with (a)p-Ga0.86In0.14N and (b) p-GaN optical gates.
www.pss-rapid.com

RapidResearch Letter

Normarized photosensitivity

Phys. Status Solidi RRL 7, No. 3 (2013)

1.0

217

cal gate. In a previous report, cutoff frequency of junction
HFET using p-GaInN has been reported to reach more than
6 GHz [15]. Thus, we consider that it is possible to further
shorten the transit time of electrons under the optical gate.
Moreover, since the minority carrier lifetime in p-GaN has
been reported to be 0.3 ns [16], photocarrier recombination
time in p-GaInN estimated to be less than 1 ns. Therefore,
in addition to the high sensitivity, high-speed response canbe expected. Dynamic performance of this device will be
examined in the near future.

Photosensor (p-GaInN)
Photosensor (p-GaN)

0.8
0.6

VAC = 5 V
RT

0.4
0.2
0.0
300

350

400

450

500

550

Wavelength [nm]Figure 3 (online colour at: www.pss-rapid.com) Normalized
monochromatic photosensitivity of HFET-type photosensors.

Photosensitivity [A/W]

were approximately 420 nm and 375 nm, respectively.
From this result, we concluded that the wavelength rangeof photosensitivity of the HFET-type photosensor can becontrolled by controlling the InN molar fraction in the
p-GaInN optical gate. In addition, the photosensor with the
p-Ga0.86In0.14N optical gate exhibits photosensitivity up to alonger wavelength of 450 nm. This is considered to be dueto the fluctuation in the InN molar fraction of p-GaInN
since this device has a thick GaInN film [13].
We also compared the photosensitivity of our photosensors with those of commercially available Si pin photodiodes (Hamamatsu Photonics K.K.: S2387-1010R) [14],Si avalanche photodiodes (Si APD) (Hamamatsu Photonics
K.K.: S8664-55) [14], and photomultiplier tubes (Hamamatsu Photonics K.K.: H7421-40, H7421-50, R928, andR3896) [14]. Figure 4 shows a summary of the monochromatic photosensitivity of each photosensor. The photosensitivity of the HFET-type photosensor with a pGaInN optical gate greatly surpassed those of the Si pin
photodiode and Si APD and was comparable to those of
the photomultiplier tubes.
Moreover, the response time of this device is thought
to be determined by transit time of 2DEG under the opticalgate or photocarrier recombination time in p-GaInN opti10

6

10

5

10

4

10

3

10

2

10

1

10

0

10

Photomultiplier tube (R3896)

Photomultiplier tube (R928)

HFET-type Photosensor(p-GaInN)
Photomultiplier tube (H7421-50)
Photomultiplier tube (H7421-40)
Si APD (S8644-55)

HFET-type
Photosensor(p-GaN)Si pin photodiode(S2387-1010R)

-1

300

450

600

750

900

Wavelength [nm]Figure 4 (online colour at: www.pss-rapid.com) Monochromatic
photosensitivity of commercially Si pin photodiode, Si APD,
photomultiplier tubes and HFET photosensors of this work.
www.pss-rapid.com

4 Summary We fabricated high-performance HFETtype photosensors for the detection of visible light. Bychanging the InN molar fraction of the p-GaInN opticalgate, the wavelength range of photosensitivity of the
HFET-type photosensors can be controlled. The photosensitivity of the HFET-type photosensors with a p-GaInN optical gate greatly surpassed those of the commerciallyavailable Si pin photodiode and Si APD and was comparable to those in photomultiplier tube.
Acknowledgements This study was supported by the
Ministry of Education, Culture, Sports, Science and Technologyof Japan, through the Supported Program for the Strategic Research Foundation at Private Universities, 2012–2016, and theKnowledge Cluster Initiative (Second Stage) of Tokai Region
Nanotechnology Manufacturing Cluster.

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© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatuspss

physicaPhys. Status Solidi C, 9, No. 3–4, 519–522 (2012) / DOI 10.1002/pssc.201100499

cwww.pss-c.com

current topics in solid state physics

Growth of GaN and AlGaN
on (100) β-Ga2O3 substrates

Shun Ito*,1, Kenichiro Takeda1, Kengo Nagata1, Hiroki Aoshima1, Kosuke Takehara1, Motoaki Iwaya1,
Tetsuya Takeuchi1, Satoshi Kamiyama1, Isamu Akasaki1,3, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Department of Electrical Engineering and Computer Science, Nagoya University, Nagoya 464-8603, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received 1 August 2011, revised 8 November 2011, accepted 8 November 2011
Published online 25 January 2012
Keywords β-Ga2O3(100), thermal annealing, GaN, AlGaN, MOVPE*

Corresponding author: e-mail 103434008@ccalumni.meijo-u.ac.jp

The crystalline quality of GaN and Al0.08Ga0.92N epitaxial
layers on (100) β-Ga2O3 substrates was significantly improved by the facet-controlled growth method. The facets
were controlled by changing the nitrogen ambient thermal annealing temperature. We demonstrated the highcrystalline-quality GaN and Al0.08Ga0.92N on β-Ga2O3

substrates, which were comparable to GaN and AlGaN
on sapphire substrates using low-temperature buffer layers. This method is useful for the fabrication of verticaltype ultraviolet (UV) light-emitting diodes (LEDs) on
β-Ga2O3 substrates.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction Group III-nitride-based UV LEDs are
expected in many applications such as excitation sourcesof white LED lamps, material processing, healthcare field,
and sterilization. There have been many reports on highefficiency UV LEDs on sapphire substrates [1, 2]. However, the wall plug efficiency of UV LEDs is still lower
than that of GaInN-based blue LEDs [3]. One of the major
causes of such low-wall-plug efficiency in the UV LEDs isthe current crowding effect caused by the combination of
insulating sapphire substrates and an n-type AlGaN cladding layer with a relatively higher sheet resistance than thatof GaN. One of the best methods to solve this problem is touse a vertical conductive device structure. Thus far,
GaInN-based vertical visible and near-UV LEDs weregrown on GaN and SiC conductive substrates [4, 5]. However, there is no report on AlGaN-based vertical UV LEDs
on conductive transparent substrates. β-Ga2O3 is one of themost attractive substrates for AlGaN-based vertical UV
LEDs. It has a transparency of up to 260 nm and n-typehigh conductivity [6]. These properties can lead to a small
absorption of UV light in β-Ga2O3 and make it possible torealize vertical conductive LEDs. Therefore, β-Ga2O3 is apromising material for the substrate of vertical UV LEDs.
Several groups have reported the GaN epitaxial growth on
β-Ga2O3 and GaInN-based blue LEDs [7, 8]. Moreover,

our group has reported AlGaN with mirror surfaces on βGa2O3 using low-temperature GaN buffer layers (LT-GaN),
which is essential for UV LEDs [9]. However, the crystalline quality of these epitaxial layers was poorer than thaton sapphire substrates. Thus, high-quality AlGaN films on
β-Ga2O3 substrates have been urgently required to achieve
high-performance UV LEDs. In this study, to obtain highcrystalline-quality GaN and AlGaN epitaxial layers on(100) β-Ga2O3 substrates, GaN and AlGaN were grown using facet layers reported previously [10]. The thermal annealing temperature of β-Ga2O3 substrates was used as aparameter to control the facet formation.
2 Experimental procedure In this study, crystal
growth was performed by metal organic vapor phase epitaxy. The aluminium and gallium sources were trimethylaluminium and trimethylgallium, respectively. The nitrogen source was NH3. Moreover, the thermal annealing andgrowth of GaN and Al0.08Ga0.92N were carried out in nitrogen ambient, because β-Ga2O3 substrates were etched usinghydrogen. The GaN and Al0.08Ga0.92N growths were carriedout using LT-GaN. AlxGa1-xN (0001) films grow epitaxially on (100) β-Ga2O3 substrates with an in-plane epitaxial
relationship of AlxGa1-xN [1-100]∥β-Ga2O3[001] [11].

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

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S. Ito et al.: Growth of GaN and AlGaN on (100) β-Ga2O3 substrates

Figure 1 shows the timing charts of the growth temperatures of GaN and Al0.08Ga0.92N without facet
AlxGa1-xN layers. After performing the thermal annealing
for 3 min at 1100 °C and deposition of LT-GaN at 550 °C,
GaN and Al0.08Ga0.92N of approximately 2.0-μm-thick weregrown at 1080 °C.

550

(d)
a: 950℃
RMST=
0.30(nm)(e)
Ta: 1100℃
RMS= 0.42(nm)RMS = 2.67(nm)RMS = 14.14(nm)Ta: 800℃

RMS = 0.20(nm)

GaN
and
AlGaN

LTGaN

(c)

(a) substrate

℃

1100
1080

(b) Ta: 600℃

time

Figure 1 Timing chart of the growth temperatures of
GaN and AlGaN.

Figure 2 shows the timing charts of the growth temperatures of GaN and Al0.08Ga0.92N grown with facetAlxGa1-xN layers. The thermal annealing for 3 min anddeposition of LT-GaN at 550 °C were performed in thesame manner as for the samples without facet-AlxGa1-xN
layers. In addition, approximately 300-nm-thick facet-GaN
or Al0.08Ga0.92N was grown at 1020 and 950 °C. The
growth temperatures of the GaN and Al0.08Ga0.92N facets
were lower than those of conventional GaN and
Al0.08Ga0.92N layers to enhance island growth. Finally, GaN
and Al0.08Ga0.92N of approximately 2.0-μm-thick weregrown on facet-GaN and Al0.08Ga0.92N at 1080 °C. In thiscase, the thermal annealing temperature was changed from
600 to 1100 °C.
Samples were characterized by atomic force microscopy (AFM), scanning electron microscopy (SEM), X-ray
rocking curve (XRC), and photoluminescence (PL) measurement using a He-Cd laser (325 nm) at room temperature.

Figure 3 Plan-view AFM images of (a) (100) β-Ga2O3 substrate and (100) β-Ga2O3 substrate thermal-annealed at (b)600 °C, (c) 800 °C, (d) 950 °C, and (e) 1100 °C for 3 min.

subsequent deposition of LT-GaN at 550 °C and growth ofapproximately 300-nm-thick GaN facet layers at 1020 °C.
From these figures, the GaN islands with inclined facets
were grown, and the area of the surface covered by the
facets was controlled by adjusting the thermal annealing
temperature.
(a) Ta: 600℃(b) Ta: 800℃

10μm(c) Ta: 950℃(d) Ta: 1100℃

Growth facet layer at
low temperature
℃
Ta
1080
1020

℃
Ta
1080
950

550

LTGaN

facetGaN

GaN
550

time

LTGaN

facetAlGaN

AlGaN

time

Figure 2 Timing charts of growth temperatures of (a)GaN and (b) AlGaN using facet layers.

3 Results
3.1 Impact of thermal annealing temperature on βGa2O3 Figure 3 shows AFM images of (a) (100) β-Ga2O3
substrate and (100) β-Ga2O3 substrates annealed at (b)600 °C, (c) 800 °C, (d) 950 °C, and (e) 1100 °C for 3 min.
From these figures, the surface roughness of β-Ga2O3 increases with increasing annealing temperature. By annealing at 1100 °C, the root mean square (RMS) roughness wasincreased from 0.20 to 14 nm.

Figure 4 Plan-view SEM images of substrates thermalannealed at (a) 600 °C, (b) 800 °C, (c) 950 °C, and (d) 1100 °C
followed by deposition of low-temperature GaN buffer layer
and growth of approximately 300-nm-thick GaN layer at1020 °C.

3.3 Growth of GaN on facet layers Figure 5 shows
plan-view SEM images of approximately 2.0-μm-thick
GaN layers grown at 1080 °C on each facet-GaN sample.
The surfaces of the GaN films on thermal-annealed substrates at (a) 600 °C and (b) 800 °C were not smooth.
However the surfaces of the GaN films on samples thermal-annealed at (c) 950 °C and (d) 1100 °C were smoothand crack-free.

3.2 Fabrication of facet-GaN Figure 4 shows planview SEM images of samples after substrate annealing and© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

www.pss-c.com

Contributed
Article521

Phys. Status Solidi C 9, No. 3–4 (2012)(a) Ta: 600℃(b) Ta: 800℃

sity of GaN on (100) β-Ga2O3 substrate was successfully
decreased by almost one order of magnitude. Figure 8
shows the PL spectra of the GaN with and without the
facet-GaN layer. These results indicated that the facet-GaN
layer improves the crystalline quality and optical propertyof GaN on (100) β-Ga2O3 substrate.

10μm
Ta: 950℃(d) Ta: 1100℃

Figure 5 Plan-view SEM images of 2.0-μm-thick GaN grown
at 1080 °C on substrates thermal-annealed at (a) 600 °C, (b)800 °C, (c) 950 °C, and (d) 1100 °C followed by growth of approximately 300-nm-thick GaN facet layers at 1020 °C.

Figure 6 shows the full width at half maximum (FWHM)
of XRC of GaN. Figure 4 indicates that the facet-GaN covered the whole surface when the thermal annealing temperature was 950 °C, and the flat surfaces remain when it
was higher and lower than 950 °C. As mentioned in reference [10], the dislocations existing in the facet regions
bend and do not glide to the surface. Therefore, the crystalline quality of GaN was highest when the annealing temperature is 950 °C. The formation of facets was related tothe adequate roughness of the β-Ga2O3 substrates. Further
evaluation is necessary.

FWHM (arcsec)

2000

1500
1000
500
0

(0002)-diffraction(20-24)-diffraction
GaN
without facet-GaN

GaN-with optimized facet-GaN

Figure 7 XRC FWHM of GaN without and with facet sample
formed by thermal annealing at 950 °C.

Intensity (arb.unit)(c)

FWHM (arcsec)

2000

GaN with optimized facet-GaN

GaN without facet-GaN

350

400

450

500

550

Wavelength (nm)1600

Figure 8 PL spectra of GaN without and with facet sample
formed by thermal annealing at 950 °C.

1200
800
400
0

(0002)-diffraction(20-24)-diffraction
600

700

800

900

1000

1100
o

Annealing temperature ( C)
Figure 6 XRC FWHM of GaN for various thermal annealing
temperatures.

Figure 7 shows the comparison of XRC FWHM between the GaN without and with the facet-GaN layer bythermal annealing at 950 °C. The threading dislocation
densities of the GaN films were estimated from the FWHM
values [12]. The calculated dislocation densities of GaN
without and with the facet-GaN layer were 1.9x1010 and
2.5x109 cm-2, respectively. The threading dislocation denwww.pss-c.com

3.4 Result of Al0.08Ga0.92N The Al0.08Ga0.92N was grownunder the same growth conditions of GaN. Only the facetAl0.08Ga0.92N growth temperature was changed to 950 °C.
Figure 9 shows a plan-view SEM images of facetAl0.08Ga0.92N and an approximately 2.0-μm-thick
Al0.08Ga0.92N layer grown on facet-Al0.08Ga0.92N samples at1080 °C. Moreover, facet-Al0.08Ga0.92N made it possible toobtain a smooth surface.
Figure 10 shows the XRC FWHM of Al0.08Ga0.92N without and with the facet-Al0.08Ga0.92N layer. The calculated
dislocation densities of GaN without and with the facetGaN layer were 2.6x1010 and 4.9x109 cm-2, respectively.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

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S. Ito et al.: Growth of GaN and AlGaN on (100) β-Ga2O3 substratesa

Ga2O3 substrates is useful for high-efficiency vertical UVLEDs in the future.

b

1μm

3μm

Figure 9 Plan-view SEM images of (a) facet-Al0.08Ga0.92N and(b) approximately 2.0-μm-thick Al0.08Ga0.92N layer grown on
facet-Al0.08Ga0.92N samples at 1080 °C.

Figure 11 shows the PL spectra of the samples without
and with the facet-Al0.08Ga0.92N layer. These results showthat the facet-Al0.08Ga0.92N layer is useful for growing
high-quality Al0.08Ga0.92N on (100) β-Ga2O3 substrate.

FWHM (arcsec)

2500
2000
1500
1000
500

(0002)-diffraction(20-24)-diffraction0

Al0.08Ga0.92N
without facet-Al0.08Ga0.92N--Al0.08Ga0.92N
with optimized facet-Al0.08Ga0.92N

Intensity (arb.unit)

Figure 10 XRC FWHM of Al0.08Ga0.92N without and with facet-Al0.08Ga0.92N formed by substrate thermal annealing at
950 °C.

Al0.08Ga0.92N with optimized facet-Al0.08Ga0.92N

Al0.08Ga0.92N without facet-Al0.08Ga0.92N

350

400

450

500

Acknowledgements This work was partially supported by
a Grant-in-Aid for Scientific Research in Priority Area "Optoelectronics Frontier by Nitride Semiconductor-Ultimate Utilization of
Nitride Semiconductor Material Potential" (No. 18069011) of the
Ministry of Education, Culture, Sports, Science and Technology.

References[1] H. Tsuzuki, F. Mori, K. Takeda, T. Ichikawa, M. Iwaya, S.
Kamiyama, H. Amano, I. Akasaki, H. Yoshida, M. Kuwabara,Y. Yamashita, and H. Kan, Phys. Status Solidi A 206,
1199 (2009).
[2] C. Pernot, M. Kim, S. Fukahori, T. Inazu, T. Fujita, Y. Nagasawa, A. Hirano, M. Ippommatsu, M. Iwaya, S. Kamiyama, I. Akasaki, and H. Amano, Appl. Phys. Express 3,061004 (2010).
[3] X. A. Cao and S. D. Arthur, Appl. Phys. Lett. 85, 3971
(2004).
[4] J. Edmond, A. Abare, M. Bergman, J. Bharathan, K. L.
Bunker, D. Emerson, K. Haberern, J. Ibbetson, M. Leung, P.
Russel, and D. Slater, J. Cryst. Growth 272, 242 (2004).
[5] Y. Narukawa, M. Ichikawa, D. Sanga, M. Sano, and T. Mukai, J. Phys. D: Appl. Phys. 43, 354002 (2010).
[6] H. Aida, K. Nishiguchi, H. Takeda, N. Aota, K. Sunakawa,and Y. Yaguchi, Jpn. J. Appl. Phys. 47, L11 (2008).
[7] K. Shimamura, E. G. Villora, K. Domen, K. Yui, K. Aoki,
and N. Ichinose, Jpn. J. Appl. Phys. 44, L7 (2005).
[8] X. Zi-Li, Z. Rong, X. Chang-Tai, X. Xiang-Qian, H. Ping,
L. Bin, Z. Hong, J. Ruo-Lian, S. Yi, and Z. You-Dou, Chin.
Phys. Lett. 25, 2185 (2008).
[9] Y. Kawai, S. Ito, K. Takeda M. Iwaya, S. Kamiyama, H.
Amano, and I. Akasaki, Ext. Abst. of ISPlasma 2010,
PB-56B (2010).
[10] S. Nitta, J. Yamamoto, Y. Koyama, Y. Ban, K. Wakao, andK. Takahashi, J. Cryst. Growth 272, 438 (2004).
[11] S. Ohira, N. Suzuki, H. Minami, K. Takahashi, T. Araki,
and Y. Nanishi, Phys. Status Solidi C 4, 2306 (2007).
[12] S. R. Lee, A. M. West, A. A. Allerman, K. E. Waldrip, D.
M. Follstaedt, P. P. Provencio, and D. D. Koleske, Appl.
Phys. Lett. 86, 241904 (2005).

550

Wavelength (nm)Figure 11 PL spectra of Al0.08Ga0.92N without and with facet
sample formed by thermal annealing at 950 °C.

4 Conclusion The thermal annealing of β-Ga2O3 substrate at 950 °C led to the form ation of GaN and
Al0.08Ga0.92N having an inclined facet covering whole surfaces, resulting in high-quality GaN and Al0.08Ga0.92N on βGa2O3 substrates. The facet-controlled method with β© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

www.pss-c.com

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Relationship between misfit-dislocation formation and initial threading-dislocation density in
GaInN/GaN heterostructures

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REGULAR PAPER

Japanese Journal of Applied Physics 54, 115501 (2015)
http://dx.doi.org/10.7567/JJAP.54.115501

Relationship between misﬁt-dislocation formation and initial threading-dislocation density
in GaInN/GaN heterostructures
Motoaki Iwaya1, Taiji Yamamoto1, Daisuke Iida1, Yasunari Kondo1, Mihoko Sowa1, Hiroyuki Matsubara1,
Koji Ishihara1, Tetsuya Takeuchi1, Satoshi Kamiyama1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

2

Received July 21, 2015; accepted August 12, 2015; published online October 1, 2015
The dependence of the critical thickness for the introduction of misﬁt dislocations in a GaInN/GaN heterostructure system on the dislocation
density in the underlying GaN layer was investigated using in situ X-ray diffraction (XRD), ex situ scanning electron microscopy, and transmission
electron microscopy analyses. The critical thickness for the introduction of misﬁt dislocations in the GaInN layer was found to signiﬁcantly dependon the dislocation density in the underlying GaN layer. Notably, on the basis of the in situ XRD results, a reliable critical thickness for obtaining
misﬁt-dislocation-free growth was determined. © 2015 The Japan Society of Applied Physics

1.

Introduction

GaInN ternary alloys have bandgaps extending from lessthan 0.65 to 3.4 eV, and therefore are used as active layers
in commercially available visible-light-emitting devices.1)
There are many interesting features that render these nitride
semiconductor alloys particularly useful as eﬃcient lightemitters. In addition, these materials are also useful in highconversion-eﬃciency solar cells.2–5) Most devices employing
GaInN ﬁlms are fabricated on GaN, because the growth of
thick GaInN ﬁlms with high crystallinity on substrates otherthan GaN is very diﬃcult. Several studies were previously
conducted to elucidate the mechanism by which defects
form via strain relaxation in GaInN=GaN heterostructures,6–11) because it is necessary to minimize defect formation in order to improve device performance. However,current understanding of the critical layer thickness at which
misﬁt dislocations (MDs) are introduced in GaInN=GaN
heterostructures is insuﬃcient. For example, the reported
values for the critical thickness in GaInN=GaN heterostructure systems vary widely. In addition, the relationshipbetween the threading dislocation (TD) density in the underlying layer and critical thickness is unclear. We previously
reported the results of in situ X-ray diﬀraction (XRD)
analyses of GaInN ﬁlms during metal organic vapor phase
epitaxy (MOVPE).12,13) By comparing these results withex situ characterization data, such as transmission electron
microscopy (TEM), cathode luminescence, atomic force
microscopy, scanning electron microscopy (SEM), and
XRD reciprocal mapping results, we found that it is possible
to accurately observe the relaxation process that occurs
in GaInN by analyzing the full width at half maximum(FWHM) of the in situ GaInN XRD spectrum.
In the present study, the critical thickness of introduced
MDs in GaInN=GaN heterostructure systems as a function of
the TD density in the underlying GaN layers was investigatedusing in situ XRD and ex situ analyses. The results indicatedthat the critical thickness for the introduction of MDs in
the GaInN ﬁlms depended signiﬁcantly on the dislocation
density in the underlying GaN layers.
2.

Experimental procedure

GaInN samples were grown on c-plane sapphire andfreestanding GaN substrates by MOVPE (Taiyo Nippon

Sanso GRC-230X) with a face down 2-in. × 3-wafer system(horizontal ﬂow). Trimethylgallium, triethylgallium (TEG),
trimethylindium (TMI), and ammonia were used as theGa, In, and N source materials for the growth of GaN andGaInN. For the sapphire substrate (∼330 µm), the GaInN
samples were fabricated on an approximately 3-µm-thick
GaN template grown using a low-temperature buﬀer layer.
The TD densities of the GaN template and free standing GaN
substrate (∼330 µm) were approximately 3 × 108 cm−2 andless than 3 × 106 cm−2, respectively. After the growth of a
1-µm-thick homoepitaxial GaN layer at 1,050 °C on the GaN
substrate or template using H2 as the carrier gas at 933 hPa,each sample was cooled to 730–800 °C and the carrier gas
was changed to N2. A GaInN layer was then grown oneach homoepitaxial GaN layer. The InN mole fraction in the
GaInN ﬁlm was varied from 0.1 to 0.22 by adjusting the
growth temperature. The mole ratio of TMI to TEG during
GaInN growth was ﬁxed at approximately 0.67.
The GaInN ﬁlms with symmetric (0002) Bragg diﬀractions
were evaluated using an in situ XRD system. The MOVPE
apparatus was equipped with a Be window that served as a
viewport for the passage of X-rays, which were focused on
the sample surface using a Johansson curved crystal mirror.
Using this method, the incidence angle of X-rays could bechanged without moving the X-ray source. Moreover, thediﬀracted X-rays were detected using a one-dimensionalcharge-coupled device, allowing for the detection of thescattered X-rays without moving the substrate or detector.
With this conﬁguration, the system realized the equivalent ofa (0002) 2θ=ω scan in 1 s during the rotation of the wafers
without requiring the use of an analyzer crystal. Furthermore,
the 2θ value was calibrated on the basis of the GaN peak.
Although the resolution slightly decreased compared withthat of a conventional XRD system, the equivalent of a(0002) 2θ=ω scan at a resolution of 1 arcsec was possible.
It is also noteworthy that the tilt mosaicity of the crystals
was reﬂected in the spectrum because an analyzer crystal
was not used. Therefore, the evaluation of the FWHM
data obtained for crystalline dislocations using this in situ
XRD system provided information on both dispersion in
the lattice constant c and tilt mosaicity. The GaInN surface structures and defects close to the GaInN=GaN
heterointerface were also evaluated using SEM and TEM,
respectively.

115501-1© 2015 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 54, 115501 (2015)

M. Iwaya et al.

ﬁlm on the GaN template was considerably ably diﬀerent
from that for the ﬁlm on the GaN substrate; the hc for the
GaInN ﬁlm on the GaN substrate (100 nm) was approximately twice that for the GaInN ﬁlm on the GaN template(55 nm).
To investigate these phenomena, Ga0.87In0.13N samples
with diﬀerent thicknesses on GaN templates and substrates
were investigated using ex situ characterization techniques.
Figures 2 and 3 show the plan-view TEM images, plan-view
SEM images, and proposed growth models for the Ga0.87In0.13N samples of diﬀerent thicknesses on the GaN templates
and substrates, respectively. As can be seen in these ﬁgures,signiﬁcant diﬀerences were observed with increasing ﬁlm
thickness.
In the Ga0.87In0.13N samples on the GaN templates, many
growth pits were observed in the plan-view SEM images
from the initial growth stage. In addition, many TDs in the
GaInN sample were observed in the GaN underlying layer inthe plan-view TEM images. However, no MDs were detectedin the 40- and 80-nm-thick Ga0.87In0.13N ﬁlms on the GaN
templates in the plan-view TEM images. Note that it waspossible to observe the GaN=GaInN interface because thethickness of each TEM sample was approximately 250 nm.
The lack of any MDs in the 80-nm-thick Ga0.87In0.13N ﬁlm
on the GaN template was also conﬁrmed via cross-sectional
TEM analysis. Therefore, lattice relaxation introduced by
MDs did not occur in the Ga0.87In0.13N ﬁlms on GaN
templates with thicknesses up to 80 nm. Moreover, the originof the growth pits was found to be the TDs in the underlying
GaN layer, as reported in Refs. 7 and 15. In contrast, only
a + c-type MDs along the h1100i
direction close to theGa0.87In0.13N=GaN heterointerface, which were induced by
slips along the h1123i
direction in the f1122g
plane
16)
system, were observed in the 110-nm-thick Ga0.87In0.13N
sample. The components of the MDs were characterized via
dark-ﬁeld cross-sectional TEM analysis. Only these a + c-

1.2
Simulation
GaInN FWHM(GaN substrate)
GaInN FWHM(GaN template)

1.0

FWHM [deg.]

0.8

hc: on GaN template (100 nm)

0.6

hc: on GaN substrate (55 nm)

0.4
0.2
0.0

0

50

100 150 200 250 300
Ga0.87In0.13N thickness [nm]

350

400

Fig. 1. (Color online) Typical in situ XRD FWHMs of the Ga0.87In0.13N
ﬁlms on GaN templates and substrates.

3.

Results and discussion

Figure 1 shows typical in situ XRD FWHMs for theGa0.87In0.13N ﬁlms on the GaN templates and substrates.
Each FWHM was determined by ﬁtting a pseudo-Voigt function for each in situ XRD spectrum. The simulated FWHM
values obtained using the X’pert Epitaxy 4.0a program14)(typical X-ray analysis software) are also presented. The
FWHMs for the Ga0.87In0.13N ﬁlms decreased with increasing
ﬁlm thickness during the initial growth stage. This result
nearly matched the simulated results. However, the experimental results deviated from the simulated results withrespect to the increase in ﬁlm thickness. Inﬂection points
were also observed in the experimental data. To clarify theseinﬂection points, which are hereafter referred to as hc,
each FWHM curve was diﬀerentiated as a function of ﬁlm
thickness. The behavior of the FWHM for the Ga0.87In0.13N

On Sapphire(GaN template)

Ga0.87In0.13N thickness
40 nm

80 nm

110 nmhc

Plan-view
Plan
view TEM
images

170 nm䚾1120䚿

䚾1120䚿
䚾1100䚿
䚾0001䚿

䚾1100䚿

䚾1100䚿

500nm

䚾0001䚿

250 nm1120
1100

䚾1120䚿

500nm

䚾0001䚿

1μm

0001

1µm

Plan
view SEM
Plan-view
imagesBent TD

Growth modelof GaInN

Pits

hc:100nm

a+c MD

Pits

GaInN

GaInN
GaN

GaN

LT-buffer
SapphireLT-buffer
Sapphirea MD
a+c MD
Pits
GaInN
GaN
LT-buffer
Sapphire

Fig. 2. (Color online) Plan-view TEM images, plan-view SEM images, and proposed growth model for Ga0.87In0.13N samples of diﬀerent thicknesses on
GaN templates.

115501-2© 2015 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 54, 115501 (2015)

M. Iwaya et al.

Ga0.87In0.13N thickness

On GaN
substrate

Plan-view TEM
images

50 nm

70 nm

hc

200 nm䚾1100䚿
䚾1120䚿

䚾1120䚿

䚾1120䚿

䚾1100䚿
500nm

䚾0001䚿

䚾1100䚿
䚾0001䚿

500nm

䚾0001䚿

500nm

Plan-view SEM
imageshc: 55 nm
Growth modelof GaInN

a MD
a+c MD
Pits

a+c MD

GaInN

GaInN

GaInN

GaN sub.

GaN

GaN

Fig. 3. (Color online) Plan-view TEM images, plan-view SEM images, and proposed growth model for Ga0.87In0.13N samples of diﬀerent thicknesses on
GaN substrates.
3

10

GaInN thickness (nm)type MDs were also observed in the 170-nm-thick Ga0.87In0.13N sample, although at a higher density than that in the
110-nm-thick ﬁlm. Both a-type and a + c-type MDs wereobserved in the 250-nm-thick Ga0.87In0.13N sample. Notably,
the hc identiﬁed in the in situ XRD analysis matched theappearance of the a + c-type MDs. Finally, these trends andrelaxation processes were similar even when the InN molar
fraction in the GaInN ﬁlm was 0.17 (see below).
Notably, for the Ga0.87In0.13N samples on the GaN substrate, growth pits and MDs were not observed during theinitial growth stage up to a Ga0.87In0.13N thickness of 50 nm.
The growth pits did not appear because the TD density in the
GaN substrate was much lower. On the other hand, a + ctype MDs were observed in the 70-nm-thick Ga0.87In0.13N
sample, and a-type MDs were detected in addition to a + ctype MDs in the 200-nm-thick Ga0.87In0.13N sample. In
addition, many growth pits were seen in the plan-view SEM
image of this Ga0.87In0.13N sample. Also, the photoluminescence intensity from GaInN ﬁlms at room temperature
rapidly decreased after the introduction of MDs in bothcases on the GaN substrate and template.
These results indicated that the GaInN ﬁlms on the GaN
substrates underwent a relaxation process similar to that forthe ﬁlms on the GaN templates, even though the growth pits
appeared at the initial growth stage for the GaInN ﬁlms
on the GaN templates. Notably, the appearance of hc wasconsistent with the introduction of the a + c-type MDs inboth GaN templates and substrates. Note that the symmetric(0002) Bragg diﬀraction was measured via in situ XRD.
Because the FWHM of this diﬀraction is dependent on the tilt
component, it is reasonable for the FWHM to increase whena + c type-MDs are introduced. Thus, we concluded that hc
indicates the critical thickness at which a + c type-MDs are
introduced.
Next, the critical thickness for the introduction of a + c
type-MDs for various InN molar fractions was investigated.
Figure 4 shows the critical thickness for the introduction

Calculation (Ref. 17)
Calculation (Ref. 18)
Calculation (Ref. 11)
Calculation (Ref. 19)
hc (on GaN template)hc (on GaN substrate)

2

10

1

10

0

10
0.00 0.05 0.10 0.15 0.20 0.25 0.30
InN molar fraction in GaInN
Fig. 4. (Color online) hc of GaInN ﬁlms on GaN templates and substrates
as a function of InN molar fraction. The results of theoretical calculations
performed using data in Refs. 11, 17–19 are also shown.

of a + c type-MDs (hc) in the GaInN ﬁlms on both GaN
templates and substrates as a function of InN molar fraction.
The results of theoretical calculations obtained using the data
in Refs. 11, 17–19 are also included in the ﬁgure. It can beseen that the critical thickness for the introduction of a + c
type-MDs in the GaInN ﬁlms depended strongly on thedislocation density of the underlying GaN layer. This difference is due to the appearance of the growth pits in the caseof the GaN template. It is possible that a fraction of the strainenergy is released as a consequence of the formation of thegrowth pits, because it is thought that the strain energy
required for the introduction of the MDs in the GaN template

115501-3© 2015 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 54, 115501 (2015)

M. Iwaya et al.

increased with critical thickness. However, note that thegrowth pits generated thinner ﬁlms on the GaN templates. In
addition, because the growth pits cause signiﬁcant negativeeﬀects with respect to device performance, GaN substrates
are advantageous when fabricating devices. Furthermore, theexperimental results for the GaInN ﬁlms on the GaN substrates closely agreed with the calculated results obtainedusing the theory of Fischer et al.19) This theory considers
the interaction between dislocations, in contrast to the othertheories. Therefore, we concluded that this theory is themost applicable since GaN ﬁlms have many dislocations.
Therefore, it is preferable to use this theory when designing
GaInN=GaN heterostructure systems for device applications
based on the introduction of a + c type-MDs.
4.

Conclusions

We investigated GaInN=GaN heterostructures on underlying
GaN layers with diﬀerent dislocation densities using bothin situ XRD and ex situ characterization techniques. It wasdemonstrated that the critical thickness for the introductionof a + c type-MDs can be obtained via in situ XRD analysis.
In addition, the critical thickness for the introduction of a + c
type-MDs in GaInN was found to depend strongly on thedislocation density in the underlying GaN layer.
Acknowledgments

This study was partially supported by the Program for
the Strategic Research Foundation at Private Universities(2012–2016) supported by the Ministry of Education,
Culture, Sports, Science and Technology (MEXT), the
MEXT Grant-in-Aid for Specially Promoted Research Grant
No. 25000011, and the MEXT Grant-in-Aid for ScientiﬁcResearch A Grant No. 15H02019.

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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JH06
http://dx.doi.org/10.7567/JJAP.52.08JH06

GaInN-Based Tunnel Junctions in n–p–n Light Emitting Diodes
Mitsuru Kaga1 , Takatoshi Morita1 , Yuka Kuwano1 , Kouji Yamashita1 , Kouta Yagi1 , Motoaki Iwaya1 ,Tetsuya Takeuchi1 , Satoshi Kamiyama1 , and Isamu Akasaki1;2
1
2Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering, Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

Received October 12, 2012; accepted December 10, 2012; published online May 31, 2013
We optimized pþþ -GaInN/nþþ -GaN tunnel junctions grown on conventional light-emitting diodes, corresponding to n–p–n structures. We
investigated two dependences at the tunnel junctions, the InN mole fraction dependence and a doping dependence. The lowest voltage drop at
the reverse-biased tunnel junction was 0.68 V at 20 mA with a 3 nm pþþ -Ga0:8 In0:2 N (Mg: 1  1020 cm3 )/30 nm nþþ -GaN (Si: 4  1020 cm3 )structure. We then found that the Mg memory effect was reasonably suppressed by using the pþþ -GaInN instead of a pþþ -GaN. At the same time,
the amount of Si doping in the following nþþ -GaN should be substantially high to overcome the Mg memory effect.
# 2013 The Japan Society of Applied Physics

1. Introduction

n+-GaN contact layer
1)

The tunnel diode was invented by Esaki in 1958. While thediode shows a negative resistance under forward bias, thediode shows an ohmic characteristic, not a rectiﬁcation characteristic under reverse bias. Therefore, the reverse-biased
tunnel junction can be a low-resistive-current path from ann-layer to a p-layer, resulting in greater ﬂexibilities for current
injection. There are some devices utilizing the reverse-biased
tunnel junctions. For example, multijunction solar cells,2,3)
multijunction LEDs,4,5) and vertical cavity surface-emitting
lasers6–9) with current conﬁnement are realized.
In nitride semiconductors, one of the remaining issues isa low hole concentration, resulting in a high resistivity ofp-type layers.10,11) A typical hole concentration in p-GaN is
at most 1  1018 cm3 ,12) and the resistivity is over 1  cm,
which is 100 times higher than the resistivity of n-GaN.
These poor characteristics cause high operating voltages
and low hole injection eﬃciencies, leading to poor device
performances. This situation is even worse for larger bandgap materials, e.g., AlGaN, which are important materials
for UV regions.13,14) Therefore, if nitride semiconductor
tunnel junctions are obtained, extensive performance improvements of nitride semiconductor devices are expected.
So far, a few reports on the nitride semiconductor tunnel
junctions have been published.15–19) The voltage drop at thetunnel junction is still high, over 1 V for some current ﬂow,
and there are no systematic reports on the optimum structureof the nitride semiconductor tunnel junctions. We then focus
our attention to GaInN because GaInN has a small band gap
energy, a high hole concentration,20–23) and potentially large
polarization ﬁeld due to the piezoelectric eﬀect.24–26) In thisstudy, we investigate the InN mole fraction dependenceof pþþ -GaInN layers in the tunnel junctions. The reversebiased tunnel junctions are characterized with n–p–n LED
structures consisting of the tunnel junctions and the conventional LEDs. In addition, Mg and Si doping dependencesof the voltage drops at the tunnel junctions are investigated.
On the basis of obtained results, inﬂuences of a Mg memory
eﬀect on the tunnel junction characteristics are discussed.
2. Experimental Methods

All the samples were grown on c-plane sapphire (0001)
substrates by low-pressure metal organic vapor phase

top n-GaN (500nm)
n++-GaN (30nm)++

p -Ga1-xInNxN (3 or 7.5nm)
p-GaN (100nm)
p-AlGaN electron-blocking layer. (20nm)
GaInN/GaN MQW
bottom n-GaN(3µm)
u-GaN(3µm)
LT-GaNc-plane sapphire(0001) substrate

Fig. 1. (Color online) Schematics of the n–p–n LED structure with thetunnel junction.

epitaxy (MOVPE). For the growth of the nitride semiconductor tunnel junctions, triethylgallium (TEGa), trimethylindium (TMIn), ammonia (NH3 ), and nitrogen were usedas the Ga, In, and N sources and a carrier gas, respectively.
We also used ethyl-bis(cyclopentadienyl)magnesium (EtCp2Mg), and silane (SiH4 ) as dopant precursors for p- andn-type layers.
As shown in Fig. 1, the n–p–n structure where the tunnel
junction was grown on the conventional LED structure wasfabricated to evaluate the reverse-biased tunnel junction ina device geometry. The structure consists of a GaN lowtemperature-deposited buﬀer layer (LT-GaN),27) an undoped
GaN layer, a bottom n-GaN layer, GaInN/GaN multiple
quantum wells (MQWs), a p-AlGaN electron blocking layer,
a p-GaN layer, the tunnel junction (pþþ -Ga1x Inx N/30 nm
nþþ -GaN), a top n-GaN layer, and a nþ -GaN contact layer.
We prepared a total of seven diﬀerent samples as shown in
Table I. The ﬁrst sample series (Sample 1 to Sample 3) isfor the study of InN mole fraction dependence in the pþþ GaInN layers. The InN mole fraction was controlled by
TMIn supply on/oﬀ and growth temperature, and calibratedby X-ray diﬀraction pattern measurements. Furthermore, thethickness of the pþþ -GaInN layer was adjusted according to
the InN mole fractions to satisfy the condition of below thecritical thickness.28,29) The second sample series (Sample 3

08JH06-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JH06

M. Kaga et al.

Table I. InN mole fractions, thicknesses, and Mg concentrations of thepþþ -GaInN in the tunnel junctions. Si concentrations in nþþ -GaN are also
shown.

Sample 1
Sample 2
Sample 3

InN mole
fractionThickness(nm)

Si concentration(cm3 )

Mg concentration(cm3 )

0

7.5

3  1020

1  1020

7.5

3  1020

1  1020

3  1020

1  1020

20

0.1
0.2

3

Sample 4
Sample 5

0.2
0.2

3
3

3  103  1020

2  1020
5  1019

Sample 60.2

3

6  1020

1  1020

3

20

1  1020

Sample 7

0.2

1:5  10

to Sample 7) is for the Mg and Si doping dependences in thetunnel junctions. Doping concentration was controlled byadjusting the SiH4 and Et-Cp2Mg ﬂows, and calibrated bysecondary ion mass spectrometry (SIMS).
A speciﬁc thermal annealing for Mg activation in thesamples was carried out in the LED fabrication process.
High temperature (725  C) and a long (30 min) thermal
annealing from etched sidewalls of the p-layers underoxygen ambient were required for suﬃcient Mg activation.
Details can be found elsewhere.30) After the annealing, top
and bottom n-electrode pads (Ti/Al/Ti/Au) were evaporated simultaneously by a standard lift-oﬀ technique. Since a
low-resistivity top n-GaN layer exists at the surface unlike
the conventional LED, no conductive transparent electrodeis required. A conventional LED with a p-contact layer anda Ni/Au semitransparent electrode was also prepared tocompare it with the LEDs containing the tunnel junctions.
The Mg and Si doping depth proﬁles of the samples were
measured by SIMS. Surface morphologies of the LEDs withthe tunnel junctions were characterized by an atomic force
microscope (AFM). In addition, interfaces at the tunnel
junctions were observed by scanning transmission electron
microscope (STEM). Current–voltage (I–V ) characteristics
were then measured with 300 m2 LEDs. Emission patternsof the LEDs were also observed at the injected current of
50 mA.

Fig. 2. (Color online) Bird’s-eye-view AFM images (5  5 m2 ) of thesurface for the LEDs with the pþþ -GaInN tunnel junctions with diﬀerent
InN mole fractions: (a) 0, (b) 0.1, and (c) 0.2.

Fig. 3. Cross-sectional STEM image of the LED with the 3 nm pþþ Ga0:8 In0:2 N tunnel junction.

3. Results and Discussion
3.1 InN mole fraction dependence of pþþ -GaInN in tunnel
junctionsFigures 2(a)–2(c) show bird’s-eye-view AFM images of theLEDs with tunnel junctions having diﬀerent InN mole
fractions in pþþ -GaInN layers (Samples 1–3). The InN mole
fractions in Figs. 2(a)–2(c) are 0, 0.1, and 0.2, respectively.
Surface roughness values of these samples are lower than
0.2 nm as the root-mean-square value. This value is similar
to that of the conventional LED. In addition, Fig. 3 showsa cross-sectional STEM image of the LED with the 3 nm
pþþ -Ga0:8 In0:2 N tunnel junction (Sample 3). Even though
the tunnel junction was heavily doped (more than 1  1020
cm3 ), smooth interfaces similar to those of the MQWs wereobtained.
Figure 4 shows I–V characteristics of the three LEDs
with the tunnel junctions and the conventional LED. The
operating voltages of the LEDs with the tunnel junctions
were 6.20, 4.60, and 4.38 V at 20 mA for 0, 0.1, and 0.2 InN
mole fractions in the pþþ -GaInN layers, respectively, while

Fig. 4. (Color online) Emission patterns of the LEDs with the pþþ GaInN tunnel junctions with diﬀerent InN mole fractions: (a) 0, (b) 0.1, and(c) 0.2. The case of conventional LED is also shown for comparison.

that of the conventional LED was 3.60 V. Since the structures underneath the tunnel junctions are the same as theconventional LED structure, the diﬀerences in voltages fromthe conventional LED correspond to the voltage drops at thereverse-biased tunnel junctions. Therefore the voltage drops
at the tunnel junctions are 2.65, 1.03, and 0.78 V at 20 mA
for 0, 0.1, and 0.2 InN mole fractions in the pþþ -GaInN
layers, respectively. The results clearly indicate that thepþþ -GaInN layers contribute to the drastic reductions involtage drops at the tunnel junctions. Figure 5 shows the

08JH06-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JH06

M. Kaga et al.

Fig. 5. (Color online) I–V characteristics of the three LEDs with pþþ -GaInN tunnel junctions with diﬀerent InN mole fractions. The conventional
LED result is also plotted.

Fig. 6. (Color online) Mg and Si doping proﬁles measured by SIMS in
the LEDs with pþþ -GaInN tunnel junctions with diﬀerent InN mole
fractions.

emission patterns of the three LEDs with the tunnel
junctions and the conventional LED. The conventional
LED shows a uniform emission pattern since the current isspreading at the semitransparent electrode. Similar uniform
emission patterns were also observed from the LEDs withthe tunnel junctions. The result indicates that good current
spreading at the top n-GaN was achieved and uniform tunnel
junctions were formed.
Figure 6 shows Si and Mg doping proﬁles measured by
SIMS at the tunnel junctions of Samples 1–3. ð1{2Þ  1020
cm3 of Mg and 2  1020 cm3 of Si were conﬁrmed. Wethen found that residual Mg concentrations in the following nþþ -layers, so-called Mg memory eﬀect, were suppressed using GaInN as the Mg-doped p-layers in the tunnel
junctions. When the InN mole fraction is 0, the residual
Mg concentration in the nþþ -layer remained high, about
5  1019 cm3 . On the other hand, with the Mg-dopedpþþ -GaInN layers, the residual Mg concentration steeply
decreased to 1  1019 cm3 in the subsequent nþþ -GaN
layers. At least we can say that the existence of the In
atoms on the surface should play an important role inthe suppression of the residual Mg incorporation. Further
investigation is necessary to clarify the mechanism.

(a)

3.2 Doping concentration dependence in tunnel junction

In order to further investigate the inﬂuences of the Mg
memory eﬀect, additional LEDs with 2 and 1=2 doping
concentrations (Samples 3 to 7) were characterized.
Figures 7(a) and 7(b) show I–V characteristics of the Si
and Mg doping dependences for the LEDs with the tunnel
junctions. The operating voltages of the LEDs are greatly
inﬂuenced by the amount of Si doping. Especially in thecase of 1=2 Si doping case (Sample 7, Si: 1  1020 cm3 ),the voltage drop at the tunnel junction was markedly
increased. Note that the residual Mg concentration in thenþþ -GaN was still high especially at the interface, someamount of Si could be easily compensated, resulting in thehigh voltage drop due to the wider depletion region.
Actually, the 2 Si doping sample (Sample 6, Si: 4 1020 cm3 ) showed further improvement, indicating thata higher Si concentration within the range used in the(b)Fig. 7. (Color online) I–V characteristics of (a) Mg and (b) Si doping
dependences for the LEDs with the tunnel junctions.

experiment is better. Interestingly, the voltage drop was
almost independent of the amount of Mg doping within therange in the experiment. Several reasons are considered. Alarger Et-Cp2Mg ﬂow could result in a large amount of Mg
incorporated, but with a more severe Mg memory eﬀect at
the same time. Alternatively, the polarization ﬁeld due to thepiezoelectric eﬀect may be more dominant than the built-in

08JH06-3# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JH06

M. Kaga et al.

ﬁeld with the impurity doping. Further investigation will benecessary to clarify the reason.
Eventually, the lowest voltage drop at the tunnel junction
was 0.68 V at 20 mA obtained by optimizing the InN mole
fraction and doping concentration, at which the resistancecorresponds to 34 . This is one of the lowest values of
the nitride-based tunnel junctions reported so far. It is stillnecessary to further reduce the voltage drop for practicaluses. A couple of further optimizations, such as reduction ofhetero barrier spikes between GaInN and GaN with gradedlayers, the position of GaInN layer (currently only the pþþ side), and the proﬁle of the Si doping, are expected.

(2000) 3137.
6) M. Diagne, Y. He, H. Zhou, E. Makarona, A. V. Nurmikko, J. Han, K. E.

7)
8)
9)
10)
11)
12)
13)

4. Summary

In summary, we have demonstrated the n–p–n LED
structures with the pþþ -GaInN/nþþ -GaN tunnel junctions.
By employing the 3 nm pþþ -Ga0:8 In0:2 N/heavily Si-doped
30 nm nþþ -GaN tunnel junction, we obtained a very low
voltage drop at the tunnel junction (0.68 V at 20 mA). We
found that the Mg memory eﬀect was reasonably suppressedusing the pþþ -GaInN layer in the tunnel junctions. At thesame time, the amount of Si doping in the subsequentnþþ -GaN should be substantially high to overcome the Mg
memory eﬀect. We found that it is essential to suppress the
Mg memory eﬀect in order to obtain low-resistivity tunnel
junctions.
Acknowledgements

14)

15)
16)

17)
18)
19)
20)

This work was supported by a Grant-in-Aid for ScientiﬁcResearch (C) (No. 23560015) and the Ministry of Education,
Culture, Sports, Science, and Technology Supported
Program for the Strategic Research Foundation at Private
Universities, Japan, 2012–2016. The SIMS measurements
were carried out by courtesy of Toray Research Center, Inc.
STEM was performed by courtesy of JFE Techno-Research
Corporation.

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High quality Al0.99Ga0.01N layers on sapphire substrates grown at 1150 °C by metalorganic
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REGULAR PAPER

Japanese Journal of Applied Physics 56, 015504 (2017)https://doi.org/10.7567/JJAP.56.015504

High quality Al0.99Ga0.01N layers on sapphire substrates grown at 1150 °C
by metalorganic vapor phase epitaxy
Shota Katsuno1, Toshiki Yasuda1, Koudai Hagiwara1, Norikatsu Koide1, Motoaki Iwaya1,
Tetsuya Takeuchi1*, Satoshi Kamiyama1, Isamu Akasaki1,2, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya 464-8603, Japan
3
Institute of Materials and Systems for Sustainability, Nagoya University, Nagoya 464-8603, Japan2*E-mail: take@meijo-u.ac.jp
Received October 19, 2016; revised November 9, 2016; accepted November 10, 2016; published online December 16, 2016
We systematically investigated metalorganic vapor phase epitaxy (MOVPE) growths of AlN layers with trimethylgallium (TMGa) supply on sapphire
substrates at 1100–1250 °C. We found that Ga incorporations into the AlN layers contributed to smooth surfaces covered with step terraces at theearly stage of the Al(Ga)N growth. In addition, a GaN mole fraction leading to the smooth surfaces was found to be around 2–3% at the beginningof growth. The Ga supply during the AlN layer growth at 1150 °C provided very smooth Al0.99Ga0.01N layers on sapphire substrates.
© 2017 The Japan Society of Applied Physics

1.

Introduction

Nitride-based deep ultraviolet light-emitting diodes (DUVLEDs) are expected to be utilized in various applicationsof medical and industrial ﬁelds. In order to fabricate the
DUV-LEDs, an AlN layer on a sapphire substrate is used
as a template. Originally the AlN epitaxial growths were
performed around 1100 °C,1–3) but later on high quality AlN
templates for eﬃcient DUV-LEDs have been achieved by
growing at above 1400 °C.4–11) At the same time, some issues
have been arose related to such a high growth temperature for
AlN. First, the wafer bowing caused by the high growthtemperature, deteriorates wafer uniformity and yield.12,13)
Second, metalorganic vapor phase epitaxy (MOVPE) reactors
require specialized heaters and growth chambers for the highgrowth temperature. In addition to that, some quartz andcarbon parts in the reactor must be frequently replaced due tothe damages caused by the high growth temperature.
If high quality AlN layers can be grown at low temperature,
the above issues will be solved. So far high quality AlN layersgrown at temperatures below 1200 °C have been attemptedwith various methods,14–20) such as NH3 pulsed injection
and=or migration enhanced epitaxy (MEE).21–24) Recently, it
was also reported that a Ga supply during the AlN growth
contributed to smooth surfaces of the AlN layers.25–27) The Ga
supply during the AlN growth could enhance surface migration of surface adatoms which leads to the two dimensional
growths. However, only a single sample was investigatedin each case, and so far systematic investigations, such as
growth parameter dependences or mechanism regarding theeﬀect of Ga in the AlN layer, have not been carried out.
In this study, we systematically investigated growths of
AlN layers under low growth temperature down to 1100 °C
by supplying Ga with various Ga=III supply ratios. The
Al0.97Ga0.03N layers with smooth surfaces were obtained by
growing even at low growth temperature, 1100–1200 °C. A
correlation between the smooth surface of the AlGaN layer
and the GaN mole fraction was clariﬁed. We also evaluatedthe material qualities of the Al0.99Ga0.01N layers.
2.

Experimental methods

All the samples in this study were grown on c-plane sapphire

substrates with a nitridation process by MOVPE.28–30) The
nitridation process was carried out along the following
procedure. First, NH3 was introduced into the MOVPE
reactor when a substrate temperature was room temperature.
Then the substrate temperature was ramping up from room
temperature to 1100 °C for 10 min and was held for 1 minwith the NH3 ﬂow. Note that no thermal etching underhydrogen was done for a sapphire substrate. After thenitridation, a 50 nm AlN buﬀer layer and an 80 nm AlN
intermediate layer were grown at 1100 °C and 1100–1250 °C,
respectively. The intermediate layer was used in order toobtain better crystal quality of a main AlN layer grown on
the intermediate layer.14) While the growth rate and the V=III
ratio for the buﬀer layer were 3.0 µm=h and 3.9, those forthe intermediate layer were 0.4 µm=h and 62, respectively.
Finally a 1.8 µm main AlN layer was grown at the same
temperature of the intermediate layer (1100–1250 °C) withand without trimethylgallium (TMGa) supply. The growthrate of the main AlN layer was 1.8 µm=h. The Ga=III supply
ratio was varied from 0 to 0.3. In addition, in order toinvestigate the surface morphology evolution during the main
AlN growth, the samples with various thicknesses of themain AlN layer (from 0 to 4.3 µm) grown at 1150 °C werealso prepared.
Surface morphologies of the samples were observed byatomic force microscopy (AFM). A Ga depth proﬁle in the
4.3 µm Al(Ga)N layer with TMGa supply was measured bysecondary ion mass spectrometry (SIMS). GaN mole fractions and crystal qualities in the samples were estimated by
X-ray diﬀraction (XRD) patterns. Threading dislocations in
the 4.3 µm AlN layer with TMGa supply were observed bytransmission electron microscopy (TEM).
3.

Results and discussion

Figure 1(a) shows the 10 × 10 µm2 AFM images of the main
AlN layers grown under various growth temperatures (1100,
1150, 1200, and 1250 °C) with various Ga=III supply ratios(0, 0.1, 0.18, and 0.3). Figure 1(b) also shows the AFM
images of the main AlN layers grown at 1100 °C with small
Ga=III supply ratios (0.01–0.03). It was found in the casewithout TMGa supply that the low growth temperatures(1100 and 1150 °C) resulted in rough surfaces with small

015504-1© 2017 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 56, 015504 (2017)(a)S. Katsuno et al.

(a)

Ga/III supply ratio
0.10

0.18

0.30

GaN mole fraction [%]

0

1100Υ
4μm

4μm

Growth temperature

RMS : 9.63 [nm]4μmRMS : 5.18 [nm]RMS : 4.83 [nm]

1150Υ
4μm

4μm

RMS : 5.47 [nm]4μmRMS : 0.18 [nm]RMS : 0.44 [nm]4μmRMS : 0.55 [nm]

1200Υ
4μm

4μm

RMS : 2.03 [nm]4μmRMS : 0.38 [nm]RMS : 0.15 [nm]

3

(i)
2

(ii)

(iii)
(iv)

(vi)

(v)

1

0

0

4μm

1

2

3

4

Thickness of AlGaN layer [μm]RMS : 0.63 [nm]

(b)
1250Υ
4μm

4μm

RMS : 1.07 [nm]4μmRMS : 1.45 [nm]RMS : 1.11 [nm]4μmRMS : 1.58 [nm]200 nm(b)
Ga/III ratio䠖0.01

4μm

RMS : 0.88 [nm]Ga/III ratio䠖0.02

4μm

RMS : 0.50 [nm](i)Ga/III ratio䠖0.03

RMS : 8.57 [nm]200 nm

200 nm(ii) RMS : 1.98 [nm] (iii) RMS : 0.47 [nm]4μmRMS : 0.27 [nm]200 nmFig. 1. (Color online) (a) 10 × 10 µm2 AFM images of the main AlN
layers grown under various growth temperatures (1100, 1150, 1250, and1250 °C) with various Ga=III supply ratios (0, 0.1, 0.18, and 0.3), and(b) 10 × 10 µm2 AFM images of the main AlN layers grown at 1100 °C withsmall Ga=III supply ratios (0.01–0.03).

columnar crystals with hexagonal (1150 °C) or circular(1100 °C) shapes, and high growth temperatures (1200 and
1250 °C) resulted in step bunching with wide ﬂat terraces.
While no eﬀect of TMGa supply was observed on thesurfaces of the samples grown at 1250 °C, very smoothsurfaces were obtained from the samples grown at 1100,
1150, and 1200 °C with diﬀerent Ga=III supply ratios, 0.03,
0.10, and 0.18, respectively. The images showed atomicallyﬂat surfaces with the steps but almost no columnar crystals or
step bunching, which RMS values were less than 0.3 nm.
Interestingly, too large Ga=III ratios also caused roughsurfaces.
We then investigated surface morphology evolutions
during the growth of the main AlN layer up to 4.3 µm,
respectively. Figures 2(a) and 2(b) show a GaN mole fraction
proﬁle measured by SIMS in the 4.3 µm main AlN layergrown at 1150 °C with 0.10 Ga=III supply ratio, and AFM
images of the surfaces of the main AlN layers with diﬀerent
layer thicknesses from 0 to 4.3 µm. Note that the 0 µm sample
corresponded to the sample showing the AlN intermediate
layer surface [Fig. 2(b)(i)]. The SIMS result reveals that
the GaN mole fraction at the very beginning of the AlN
main layer was high, 2.1%. Then the GaN mole fraction
was decreased down to 1.2% within the ﬁrst 1.3 µm thickness,
resulting in a graded layer. After that, the GaN mole fraction
remained constant, 1.2%. In other words, the two AlGaN
layers, (the graded layer and the constant layer), wereautomatically formed even under the same growth condition.
In addition, AFM images indicated that the surfaces becamesmooth (RMS values: 8.57 to 0.13 nm) during the ﬁrst 1.3 µm
layer, which was coincident with the thickness of the graded
AlGaN layer. After the ﬁrst 1.3 µm layer, the surface
smoothness was kept to be the same up to 4.3 µm. These
results suggest that TMGa supply was very useful to form
smooth surface of the main Al(Ga)N layer.

200 nm

200 nm(iv) RMS : 0.13 [nm] (v) RMS : 0.13 [nm] (vi) RMS : 0.12 [nm]Fig. 2. (Color online) (a) a GaN mole fraction proﬁle measured by SIMSin the 4.3 µm main AlN layer grown at 1150 °C with 0.10 Ga=III supply ratio,and (b) 1 × 1 µm2 AFM images of the surfaces of the main AlN layers with
diﬀerent layer thicknesses, (i) 0 µm (the intermediate layer), (ii) 0.5 µm,(iii) 0.9 µm, (iv) 1.2 µm, (v) 2 µm, and (vi) 4.3 µm.

(a)
AlN template layer(b)

AlN without
TMGa supply (0.5μm)
Graded AlGaN with
TMGa supply (1.3μm)
AlN intermidiate layer
AlN buffer layerc-plane sapphire

4μm

RMS : 0.14 [nm]Fig. 3. (Color online) (a) the sample structure and (b) a 10 × 10 µm2 AFM
image of the sample.

In order to further investigate the eﬀect of Ga atoms forthe growth of the smooth AlGaN layer, another sample consisting of the ﬁrst 1.3 µm main graded AlGaN with TMGa
supply and the second 0.5 µm main AlN without TMGa
supply was prepared. Figures 3(a) and 3(b) show the sample
structure and the AFM image of the sample. We clearlyobserved a very ﬂat surface which was similar to that of thesample with TMGa supply during all the growth of the main
AlN layer [Fig. 1(a) 1150 °C for growth temperatureand 0.10 for Ga=III supply ratio], indicating that TMGa
supply plays an important role in the ﬁrst 1.3 µm layer for thesmooth surface.
Next, we estimated the GaN mole fractions in the main
AlN layers grown under various growth temperatures andvarious Ga=III supply ratios. Figure 4(a) shows (0002) 2θ=ω
XRD patterns of the main Al(Ga)N layers grown at 1150 °C
with various Ga=III supply ratios. While a single main peak015504-2© 2017 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 56, 015504 (2017)

S. Katsuno et al.

Growth temperature=1150

(a)

(a)
1.0

35.6

35.8

36.0

36.2

Normalized Intensity [arb.unit]

Intensity [arb.unit]

Ga/III ratio = 0
Ga/III ratio = 0.10
Ga/III ratio = 0.18
Ga/III ratio = 0.30

36.4

2theta [deg]

0.1

0.2
Ga/III ratio

0.6
0.4
0.2
0.0
-1000

-500

0

500

1000

500

1000

(b)

o

1.0
Normalized Intensity [arb.unit]

0.0

0.8

omega [s]Growth temperature 1100 Co
Growth temperature 1150 CoGrowth temperature 1200 Co
Growth temperature 1250 C

GaN mole fraction [%](b) 㻝㻝
㻝㻜
㻥
㻤
㻣
㻢
㻡
㻠
㻟
㻞
㻝
㻜

sample A
sample B

0.3

Fig. 4. (Color online) (a) (0002) 2θ=ω XRD patterns of the main AlN
layers grown at 1150 °C with various Ga=III supply ratios, and (b) GaN mole
fractions of the very bottom part in the graded AlGaN layer as a function of
Ga=III ration and growth temperatures.

was observed from the AlN layer grown without TMGa
supply, two peaks were observed from the main AlN layersgrown with TMGa supplies. From the SIMS results shown in
Fig. 2(a), the two peaks of XRD patterns should correspondto the two AlGaN layers, the ﬁrst 1.3 µm graded AlGaN layer
and the second 0.5 µm constant AlGaN layer. A peak at alarger angle (lower GaN mole fraction) should correspondto the second constant AlGaN layer and a peak at a smaller
angle (higher GaN mole fraction) should correspond to thepart of the highest GaN mole fraction in the ﬁrst graded
AlGaN layer, that is the very bottom part of the graded
AlGaN layer. Furthermore, with an increase of the Ga=III
supply ratios, the two peaks were shifted to a lower angle
side, meaning that larger Ga incorporation was obtained withan increase of the Ga=III supply ratio. We then plotted the
GaN mole fractions of the very bottom part in the graded
AlGaN layer as a function of Ga=III ratio and growthtemperatures in Fig. 4(b). Coincidently the three smoothsurfaces observed by AFM were obtained in the case that theGaN mole fractions were within 2–3% (indicated with solid
circles) regardless of the growth conditions. In other words,
the results implies that not growth conditions but a certain
GaN mole fraction value seems key factor to obtain smoothsurfaces in the case of adding Ga in the AlN layer.
Finally we investigated material qualities of AlGaN mainlayers. Figures 5(a) and 5(b) show XRD ω scans of ð2024Þand ð1015Þ planes of two diﬀerent samples, respectively.
Sample A is the sample grown at 1150 °C, consisting of the

Sample A
Sample B

0.8
0.6
0.4
0.2
0.0
-1000

-500

0
omega [s]

 and (b) ð1015Þ
 planes
Fig. 5. (Color online) XRD ω scans of (a) ð2024Þ
of two diﬀerent samples.

ﬁrst 1.3 µm graded AlGaN with TMGa supply and the second
0.5 µm main AlN without TMGa supply shown in Fig. 3.
Sample B consisted of the 1.8 µm main AlN layer grown at
1250 °C without Ga supply shown in Fig. 1(a). Two peaks
were clearly observed in Sample A due to the existenceof the two layers, the 1.3 µm graded AlGaN and the 0.5 µm
AlN, and the narrower peak on the right side corresponds tothe 0.5 µm AlN. The FWHMs of the peak from the AlN layer
in Sample A are narrower than those of the peaks obtained
from the Sample B. Figure 6 shows a cross-sectional TEM
image of the 4.3 µm AlN layer with TMGa supply (the sameas the SIMS sample in Fig. 2), taken with the reﬂectionof g ¼ ½1120.
In the image, pure edge dislocations andedge=screw-mixed dislocations are observed, and the estimated dislocation density was approximately 5 × 109 cm−2,
which is comparable to that of the AlN layer grown at
1300 °C or higher.31) These results clearly suggest that Ga
supply during the AlN layer growth at less than 1200 °C isvery useful to obtain reasonably high quality AlN layer onsapphire.
4.

Conclusions

We investigated an epitaxial growth of AlN layer with Ga
supply on sapphire at relatively low growth temperature.
Smooth surfaces of the AlGaN layers with the small amountof GaN mole fractions were obtained even with the low
growth temperature, such as 1100–1150 °C. Such smoothsurfaces have been formed during the initial growth of the015504-3© 2017 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 56, 015504 (2017)

g=[ 1 1 -2 0 ]S. Katsuno et al.

1μm

Fig. 6. A cross-sectional TEM image of the 4.3 µm AlN layer with TMGa
supply (the same as the SIMS sample), taken with the reﬂections ofg ¼ ½1120.

AlN ﬁlmy layer with the TMGa supply. The Ga incorporation
could contribute to the promotion of lateral growth in theinitial growth. We also found that the samples with smoothsurfaces contained 2–3% GaN mole fractions regardless ofthe growth conditions.
Acknowledgments

This work was supported by Technology Supported Program
for the Strategic Research Foundation at Private Universities(2012–2016), MEXT Grant-in-Aid for Specially Promoted
Research (No. 25000011), and Grant-in-Aid for ScientiﬁcResearch (B) (No. 26286045).

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REGULAR PAPER

Japanese Journal of Applied Physics 55, 05FD01 (2016)
http://doi.org/10.7567/JJAP.55.05FD01

GaNSb alloys grown with H2 and N2 carrier gases
Daisuke Komori1*, Kaku Takarabe1, Tetsuya Takeuchi1, Takao Miyajima1,
Satoshi Kamiyama1, Motoaki Iwaya1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan2*E-mail: 143434022@ccalumni.meijo-u.ac.jp
Received November 5, 2015; accepted November 8, 2015; published online February 26, 2016
We grew GaNSb layers with H2 and N2 carrier gases by metalorganic vapor phase epitaxy. We estimated the GaSb molar fraction in a GaNSb layergrown with H2 by Rutherford backscattering spectroscopy. A 0.8% GaSb molar fraction was obtained, which was consistent with the value obtainedby secondary ion mass spectroscopy. We correlated the obtained GaSb molar fraction with the c-axis lattice constant of GaSb estimated from an
X-ray diffraction pattern. We investigated GaSb molar fractions in GaNSb grown with H2 and N2 at various growth temperatures. While GaSb molar
fractions in the H2 case showed a plateau at 0.8% at less than 800 °C, those in the N2 case increased to 1.1% with a decrease in the growthtemperature to 750 °C. Sb incorporation into GaNSb could be further improved by carrying out growth under N2, similar to the case of GaInN
growth. © 2016 The Japan Society of Applied Physics

1.

Introduction

Group-III nitride materials, such as AlN, GaN, InN, andrelated alloys, have been extensively investigated towards
their use in various applications, such as optoelectronics andpower electronics. For instance, light-emitting and lightdetecting devices covering near-infrared, visible, and ultraviolet regions are expected since the materials have widebandgap regions from 0.7 to 6.0 eV. Unfortunately, thematerial choices for high-quality nitride-based heterostructures have been limited because of large mismatches in thegrowth temperatures and lattice constants. Whereas a temperature of around 1400 °C is necessary for AlN growth bymetalorganic vapor phase epitaxy (MOVPE),1–4) In-containing nitride-based materials must be grown at less than 800 °C
to enable reasonable amounts of In to be incorporated.5–8)
Such a huge temperature diﬀerence results in poor interface
qualities of the heterostructures. Also, lattice mismatches for
AlN=GaN and InN=GaN are 2.4 and 11%, respectively.9,10)
These values are too large to fabricate high-quality heterostructures without the generation of dislocations.
Under the circumstances, the use of GaNSb alloys could beconsidered to solve the above challenging issues. Antimony(Sb) is known as a surfactant that enhances the migrationof surface species and increases the critical thickness ofheteroepitaxy in several semiconductor materials. For
instance, GaIn(N)As multiple quantum wells (MQWs) grown
on GaAs with Sb supply showed larger critical thicknesses
and better optical qualities.11) Furthermore, GaN layers and
GaInN MQWs grown with Sb supply showed an enhancement in the lateral growth and better optical qualities,
respectively.12,13) Thus, supplying Sb could be eﬀective forgrowing high-quality nitride-based materials at low growthtemperatures. Also, since AlGaInSb alloys have large lattice
constants, AlGaInNSb alloys could be new candidates forlattice-matched heterostructures.
At the same time, group-III nitride alloys are generally
formed by replacing the group-III elements, but research on
nitride-based alloys replaced with other group-V elements isvery limited.14–17) Regarding GaSb molar fractions in GaNSb
alloys, a 4% GaSb molar fraction in a GaNSb layer grown by
molecular beam epitaxy was estimated by Rutherford backscattering spectrometry (RBS).18) On the other hand, a 0.05%

GaSb molar fraction in a MOVPE-grown sample wasestimated by secondary ion mass spectroscopy (SIMS)
calibrated with an Sb-implanted Si standard sample.19) Thus,a more precise evaluation of MOVPE-grown GaNSb isnecessary.
In this study, we estimated the GaSb molar fraction in a
MOVPE-grown GaNSb layer by Rutherford backscattering
spectroscopy. Then, we correlated the molar fraction with thec-axis lattice constant of the same sample measured from Xray diﬀraction (XRD) patterns. We systematically investigated the carrier gas dependence, H2 or N2, of GaSb molar
fractions in GaNSb layers grown at various temperatures. We
found that the use of N2 carrier gas led to some improvements
in Sb incorporations into GaNSb layers, similar to the causeof GaInN. We also found that a Sb supply is eﬀective forobtaining smooth surfaces even at a low growth temperature
down to 750 °C.
2.

Experimental procedure

All the samples in this study were grown on 3 µm GaN
templates prepared on c-plane sapphire substrates with the
buﬀer layers deposited at low temperatures by MOVPE.20)
Then, GaNSb layers were grown on the GaN templates by
MOVPE with an Sb source. Trimethylgallium (TMGa), trimethylantimony (TMSb), and ammonia (NH3) were used as
Ga, Sb, and N sources, respectively. During all the GaNSb
epitaxial growths, supply rates of TMGa, TMSb, and NH3
were ﬁxed at 155 µmol=min, 548 µmol=min, and 152
mmol=min, respectively, resulting in a N=Ga ratio of 980
and an Sb=N ratio of 0.0036. In contrast, the carrier gas andgrowth temperature for GaNSb growth were changed. H2 or
N2 carrier gas was used during the GaNSb growth. Also,
the growth temperature was changed from 650 to 950 °C. The
thicknesses of the GaNSb layer were 500 and 200 nm with H2and N2 carrier gases, respectively. This diﬀerence was causedby the growth rate diﬀerence between H2 and N2 carrier gases.
We then measured the number of Sb atoms in the samples
by RBS and SIMS. XRD patterns of the samples were alsomeasured. Furthermore, images of the sample surfaces wereobtained by atomic force microscopy (AFM).
3.

Results and discussion

The composition of the GaNSb=GaN sample was measured

05FD01-1© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD01 (2016)

D. Komori et al.

Fig. 1. (Color online) Depth proﬁles of Sb and Ga compositions in
GaNSb=GaN measured by RBS. The GaNSb layer was grown at 800 °C with
H2.

by RBS. For the measurement, a 500 nm GaNSb layer
was grown on the 3 µm GaN template with H2 carrier gas
at 800 °C. Figure 1 shows the measured Ga, N, and Sb
compositions in the GaNSb=GaN sample. Note that C and O
compositions were less than the detection limits in both the
GaNSb and GaN layers. The Ga, N, and Sb compositions in
the GaNSb layer were 0.500 ± 0.005, 0.496 ± 0.025, and
0.004 ± 0.001, respectively. In the bottom GaN layer, Ga, N,
and Sb compositions were 0.500 ± 0.005, 0.500 ± 0.025,
and 0.000 ± 0.001, respectively. While the Ga composition
was consistently 0.500 in both the GaN and GaNSb layers,
the N composition was reduced from 0.500 in GaN to 0.496
in GaNSb. At the same time, the Sb composition wascomplementarily increased from 0 in GaN to 0.004 in
GaNSb. Therefore, it is most likely that Sb atoms were
replaced with N atoms in GaNSb. On the basis of theseﬁndings, the GaSb molar fraction in GaNSb was estimated tobe 0.008 ± 0.002, in other words, 0.8%.
SIMS measurement was also performed on the same
GaNSb layer. The Sb concentration was found to be 3.4 ×
1021 atoms=cm3, which corresponds to a 0.9% GaSb molar
fraction. The value was calibrated with an Sb-implanted GaN
standard sample and determined to be consistent with thatobtained by RBS.
We then measured two XRD patterns of the same GaNSb
sample. Figures 2(a) and 2(b) show a (0002) 2θ=ω XRD
 reciprocal space mapping for GaNSb,
pattern and a (2024)
respectively. A clear additional peak on the left side of the
GaN peak was observed. Since GaSb has a large lattice
constant, GaNSb should have a larger lattice constant than
GaN, leading to a smaller diﬀraction angle of (0002) than thatof GaN. Thus, this peak is assigned as a GaNSb peak. In
addition, it is also found that the GaNSb layer was coherently
grown on the bottom GaN template and still fully strained,as shown in Fig. 2(b). Thus, we determined that the c-axis
lattice constant of the fully strained GaNSb with a 0.8% GaSb
molar fraction on GaN was 5.196 Å. Also, the c-axis lattice
constant of GaN was 5.189 Å. Note that a weak peak wasalso found on the right side of the GaN peak in Fig. 2(a). At
the same time, no corresponding peak was found in Fig. 2(b).
At this moment, the origin of the peak is unknown andfurther investigations are necessary. Assuming that there is alinear relationship between the GaSb molar fraction and thec-axis lattice constant, we deriveLc ¼ 5:189 þ 0:875x:

ð1Þ

 RSM
Fig. 2. (Color online) (a) (0002) 2θ=ω XRD pattern and (b) (2024)of the GaNSb sample grown at 800 °C with H2.

Fig. 3. (Color online) (0002) 2θ=ω XRD patterns of the GaNSb=GaN
samples grown at various temperatures with (a) H2 and (b) N2.

Here, x is the GaSb molar fraction and Lc is the c-axis lattice
constant of GaN1−xSbx. We later use the equation to estimate
GaSb molar fractions of GaNSb samples grown undervarious growth conditions.
Figures 3(a) and 3(b) show (0002) 2θ=ω XRD patterns of
the GaNSb=GaN samples grown at various growth temperatures with H2 and N2 carrier gases, respectively. GaNSb
peaks were clearly observed on the left side of the GaN peaksin the samples grown with H2 carrier gas at 750 °C and higher.

05FD01-2© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD01 (2016)

D. Komori et al.

growth temperature (800 °C), the surfaces were surprisingly
smooth. Also, small grains were found on GaNSb grown with
N2 carrier gas. It seems that Sb acts as a surfactant on the
GaN surface. Thus, further optimization of the growth
conditions of nitride-based materials with Sb supply could
lead to the possibility of high-quality nitride-based materials
and heterostructures being grown at low temperatures.
4.

Fig. 4. (Color online) GaSb molar fractions estimated from c-axis lattice
constants of GaNSb layers grown with H2 and N2, as a function of growthtemperature.

Conclusions

We have grown GaNSb layers with H2 and N2 carrier gasesby MOVPE. A 0.8% GaSb molar fraction in one of the
GaNSb layers was obtained by RBS. We correlated theobtained GaSb molar fraction with the c-axis lattice constantof GaNSb estimated from an X-ray diﬀraction (XRD) pattern.
The maximum GaSb mole fraction in this study was 1.1%
in GaNSb grown with N2 carrier gas at 750 °C. Also, a
relatively smooth surface of the GaNSb layer grown with
H2 was obtained, probably owing to an Sb surfactant eﬀect.
The use of GaNSb could open up the possibility of growing
high-quality nitride-based heterostructures at relatively low
temperatures.
Acknowledgements

Fig. 5. (Color online) AFM images of GaNSb surfaces grown at 800 °C
with (a) H2 and (b) N2.

No peak other than the GaN peak was observed in the sample
grown at 650 °C, meaning that the GaNSb quality was very
poor. On the contrary, broad peaks related to GaNSb wereobserved in the samples grown under N2 carrier gas. Evenconsidering the diﬀerence in the GaNSb thickness between
the cases of H2 and N2, the quality of GaNSb grown with N2
seemed to be poorer than that of GaNSb grown with H2.
Next, we estimated GaSb molar fractions from themeasured c-axis lattice constants using the above-mentioned
equation and plotted the results in the cases of H2 and N2 as
a function of growth temperature in Fig. 4. The GaSb molar
fraction in the case of 650 °C with H2 was measured by
SIMS. Basically, the GaSb molar fractions with H2 and N2
similarly increased with a decrease in growth temperaturefrom 950 to 800 °C. This was probably due to lower Sbdesorption from the growth surface at such low growthtemperatures. However, the GaSb molar fraction under H2carrier gas became saturated at about 0.8% as the growthtemperature was further decreased from 800 to 650 °C. On
the contrary, the GaSb molar fraction with N2 was monotonically increased to 1.1% at 750 °C. This value is the highest
GaSb molar fraction obtained in this study. The result
indicates that using N2 carrier gas in GaNSb growth seems tobe eﬀective for obtaining higher GaSb molar fractions.
Finally, Fig. 5 shows AFM images of GaNSb surfaces
grown at 800 °C with (a) H2 and (b) N2 carrier gases.
Typically, many pits are observed on the GaN surface grown
at 800 °C owing to insuﬃcient migration. Here, signiﬁcantdiﬀerences appeared in the surface morphologies of the
GaNSb layers. No pits but many hexagonal regions withspiral steps were covered on the GaNSb surfaces, resulting
in relatively smooth surfaces. Actually, considering the low

This study was supported by the MEXT Program for theStrategic Research Foundation at Private Universities(2012–2016) and the Grant-in-Aid for Challenging Exploratory Research (No. 15K13959), and the Foundation for
Promotion of Material Science and Technology of Japan(MST Foundation). Sb sources were provided by Japan
Advanced Chemicals.

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05FD01-3© 2016 The Japan Society of Applied Physics

Advantages of the moth-eye
patterned sapphire substrate for thehigh performance nitride based LEDs

solidistatuspss

physicaPhys. Status Solidi C 11, No. 3–4, 771– 774 (2014) / DOI 10.1002/pssc.201300492

cwww.pss-c.com

current topics in solid state physics

Toshiyuki Kondo*,1,2, Tsukasa Kitano1, Atsushi Suzuki1, Midori Mori1, Koichi Naniwae1,
Satoshi Kamiyama1,2 Motoaki Iwaya2, Tetsuya Takeuchi2, and Isamu Akasaki2
1

EL-SEED Corp., Innovative Science and Technology Building 2F, Meijo University, 2-1522 Shiogamaguchi, Tenpaku-ku,
Nagoya 468-0073, Japan
2
Faculty of Science and Technology, Meijo University, 1-501 Shiogamaguchi, Tenpaku-ku, Nagoya 468-8502, Japan
Received 23 August 2013, revised 23 December 2013, accepted 20 January 2014
Published online 21 March 2014
Keywords moth-eye structure, sapphire, LED, nitride, light extraction, PSS* Corresponding author: e-mail toshiyuki_kondo@elseed.com

We propose moth-eye patterned sapphire substrate(MPSS) as a solution to improve the light extraction efficiency and to reduce the production cost of LEDs. The
MPSS samples’ surfaces consisted of a triangular grid of
about 375 nm high truncated cones with different pitchesof 460 nm, 500 nm, 600 nm, 700 nm and 800 nm. A
commercially available patterned sapphire substrate(PSS) and a flat sapphire substrate (FSS) was also used in
the experiments for comparison. According to the cathode luminescence observation, the GaN template on
MPSS with a thickness of 3 µm showed a threading dislocation density (TDD) of around 1.9 x 108 cm-2. Transmission electron microscope observation revealed thatmany of the dislocations were bent and disappeared asloops formed in the vicinity of MPSS cones. On the other
hand, PSS required a GaN template thicker than 5µm toachieve a level of TDD equal to MPSS. The LED on
MPSS with a pitch of 600 nm showed the highest light
output power among the evaluated LEDs as it was 1.89
times higher than that on FSS and 1.05 times higher thanthat on PSS. The pitch dependence of the light output
improvement agrees with the pitch dependence of thesimulated transmissivity at the GaN/sapphire interface.
As a result of our comparison, we concluded that MPSS
provides the most cost effective solution for high performance LEDs.
© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction
General illuminations accounts for approximately 16%of all electricity consumption in households. It is expectedthat energy saving rates will increase by the replacement ofexisting illumination solutions with next-generation illumination, light emitting diodes (LED) and organic electroluminescence (OEL) that have high energy efficiency. This
technique is necessary to make sure that the nextgeneration illumination solutions obtain about twice theluminous efficiency of current fluorescent lamps and to get
the price reduction that is expected to be attained throughthe general technology development of next-generation illumination.
LED performance depends on internal quantum efficiency (IQE) and light extraction efficiency (LEE). As forthe blue LEDs, a sufficiently high IQE has already beenachieved. On the other hand, LEE is limited by the total in-

ternal reflection, which is caused by the large difference inthe refractive indexes of the group III nitride based epitaxial film and the sapphire substrate. Patterned sapphire
substrates (PSS) using micro-scale patterns and thin film
technology are generally being used in order to address thisproblem. This technology is however plagued by wafer
bowing problems that limit the possibility to increase substrate size and suffers from high production costs.
We have proposed that a new approach based on a submicron scale patterned sapphire substrate, the so-called“moth-eye patterned sapphire substrate” (“MPSS”), should
be used instead. It is a very promising solution which improves the LEE of the nitride-based LEDs [1-6].
MPSS, with its sub-micron scale structure, can, in contrast to the PSS, be successfully planarized even when one
uses a relatively thin GaN epitaxial layer.

© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

772

solidistatusphysicapsscT. Kondo et al.: Moth-eye patterned sapphire substrate for high performance nitride based LEDs

In this paper, we will make a comprehensive comparisonof LED performance and GaN layer quality on MPSS, PSS
and flat sapphire substrate (FSS).
2 Experimental
The MPSS samples were fabricated by using nanoimprinting lithography and ICP dry etching. The nanoimprinting lithography can pattern wafers with a diameterof up to 6 inches in size. The surfaces of the samples usedin our experiment consisted of a triangular grid of approximately 375 nm-high truncated cones with different
pitches: 460 nm, 500 nm, 600 nm, 700 nm and 800 nm respectively. The top of the truncated cones were rounded by
hot phosphoric acid etching to improve the crystalline
quality of the GaN layers. A bird’s-eye view SEM imageof an MPSS sample with a 460 nm pitch is shown in Fig. 1.
Commercially available PSS samples with a pitch of 3 μm
and 1.5 μm high pointed cones as well as FSS were also
used in the experiments for comparison.

Figure 1 Bird’s-view SEM image of a MPSS with a pitch of 460
nm.

Metal orgnic vapor phase epitaxy (MOVPE) was used forthe nitride layer growth and trimethylgallium, trimethylindium, and ammonia were used as source materials. Furthermore, magnesium and tetra methyl silane were used as
p-type and n-type dopant sources respectively. Prior to
growth, all of the MPSS, PSS and FSS samples were thermally cleaned at 1100 °C in hydrogen gas to get rid of surface contamination.
The GaN templates which consisted of a low temperature
GaN buffer layer and high-temperature undoped GaN layer
were grown on MPSS, PSS and FSS so that we could evaluate the crystalline quality of the GaN layer on each substrate. The thickness of the GaN template on MPSS and
FSS was 3 μm compared to 5 μm on the PSS. The LED
structures were also grown on the templates that consistedof 3 μm thick n-type GaN:Si layer, and a 0.15 μm thick ptype GaN layer. The MQW active layer consists of fivepairs of 3 nm thick GaInN wells and 12 nm thick GaN barriers, respectively. Flip-chip type blue LEDs were fabricat-© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

ed with an area of 600 µm x 500 µm. A combination of an
ITO and an Ag-Pd-Cu alloy (APC) was employed as a pelectrode and an n-electrode in order to obtain both low
contact resistance and high reflectivity [4, 9].
The LED chips were mounted onto the header without
resin encapsulation. The crystalline quality of the GaN
templates was characterized using X-ray diffraction (XRD),
cathode luminescence (CL) and transmission electron microscopy (TEM). The light output versus current (L-I) andthe angular intensity distribution of the LEDs were evaluated at room temperature (RT). In addition, we simulated
the transmissivity at the GaN/sapphire interface as a function of the incidence angle using rigorous coupled wave
analysis (RCWA).
3 Results and discussion
The full width at half maximums (FWHMs) of the XRD
for GaN(0002) and GaN(10-12) on each substrate were as
follows: 236 arcsec and 320 arcsec in case of having a
MPSS with a pitch of 460 nm; 192 arcsec and 290 arcsec
in case of having a PSS; 216 arcsec and 314 arcsec in caseof having a FSS. The thickness of the GaN template on the
MPSS (3 μm) was thinner than that on the PSS and equalin thickness to that on the FSS, yet did not cause any adversary effects on the crystalline quality nor the surface
morphology of the GaN. Having a thinner GaN layer reduces wafer bowing, which opens up for the possibility toenlarge the size of the wafer. Consequently, wafer bowing
decreased by approximately 10% as we reduced the GaN
thickness from 5 µm to 3 µm when growing GaN on sapphire substrates with a diameter of 4 inches. In addition,
the epi cost can also be reduced by growing thinner GaN
layers. Furthermore, Fig. 2 shows the CL results after GaN
template growth on MPSS with a pitch of 460 nm and on
PSS, respectively.

Figure 2 Plan view CL image of GaN template on MPSS with apitch of 460 nm (a) and that on PSS (b).

The threading dislocation density (TDD) in the GaN templates on MPSS and PSS were 1.9 x 108 cm-2 and 1.58 x
108 cm-2, respectively. In addition, the TDD of the GaN
templates on MPSS with a pitch of 500 nm, 600 nm and
800 nm were 1.9 x 108 cm-2, 2.0 x 108 cm-2 and 2.1 x 108
cm-2, respectively. However, if the bottom diameter of the
MPSS cones with a pitch of 460 nm exceeds 360 nm, theGaN surface was getting roughened due to the difficulty togrow cores on c-plane areas. The crystalline quality of
GaN templates on PSS was slightly better than that on
MPSS and FSS substrates. However, it was necessary towww.pss-c.com

Contributed
ArticlePhys. Status Solidi C 11, No. 3–4 (2014)

grow more than 5 µm thick GaN layers on PSS in order toplanarize the surface and keep a low dislocation density.
Figure 3 shows a cross-sectional bright-field TEM imageof the GaN template on MPSS with a pitch of 460 nm.
Many of the dislocations were bent and disappeared asloops formed in the vicinity of MPSS cones. These observations clearly suggest that the MPSS can provide a quality
equivalent to that of GaN templates with a thinner GaNlayer compared to that on the PSS.

773

The distinct total internal reflection of the light entering into the moth-eye plane disappeared when the incident angle
got higher than 46 degrees. The MPSS do however get alower transmissivity within the critical angle, compared to
the FSS and PSS since we get a diffractive reflectance at
the GaN/MPSS interface. The transmissivity does, however, increase due to the diffraction effect we get at angles
above the critical angle.

θ

ds

dθ

A

B

Figure 5 A schematic hemisphere used to estimate light extraction efficiency, LEE.

The following Eq. (3) was used to calculate the overall
transmissivity for each angle between the angles 0 degrees
to 89 degrees. Each variables are described in Fig. 5.
ds = A ◊ B = 2π r 2 ◊ sinθ ◊ dθ ,

(1)

0D

S = Ú D 2π r 2 ◊ sinθ ◊ dθ = 2π r 2 ,90

Figure 3 Cross sectional TEM image of the GaN template on
MPSS with a pitch of 460 nm.

MPSS has an optical wavelength order periodic structure
on a sapphire substrate, utilizing the Bragg diffraction effect to not only suppress the occurrence of total reflectionin the GaN/sapphire interface, but also extract light with anincident angle larger than the critical angle. Figure 4 shows
the transmissivity at the GaN/sapphire interface as a function of the incidence angle simulated by using RCWA [1, 7,
8].
For this simulation we assumed that the cone pattern of the
MPSS consisted of a periodic triangular grid consisting of
fixed truncated cones with a height of 375 nm, a bottomdiameter of 350 nm and a top-diameter of 150 nm. We
used different cone pitch parameters.

T = 1/ 2π r 2

θ2

(2)

Úθ T (θ ) ◊ 2π r ◊ sinθ ◊ dθ
2

1

θ2

= Ú T (θ ) ◊ sinθ ◊ dθ .
θ1

(3)

r is the radius of a sphere, T is the transmissivity of the incident angle θ calculated by simulation with RCWA.
The integrated transmissivity increased at incident angles
higher than the critical angle due to Bragg’s diffraction effect in cones with a pitch of up to 600 nm as shown in Fig.
6. On the other hand, the integrated transmissivity starts todecrease as the pitch reaches above 600 nm. This occurs
due to the increase of internal total reflection intensity as
the flat areas between the cones widen along with the increasing pitch.

Figure 4 Calculation of the transmissivity dependence on incident at the GaN/sapphire interface.
www.pss-c.com© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

774

solidistatusphysicapsscT. Kondo et al.: Moth-eye patterned sapphire substrate for high performance nitride based LEDs

GaN/sapphire interface whereas the experiment results
show the impact on LEE.

Figure 6 Calculated transmissivity over the entire solid angle atthe GaN/sapphire interface as a function of the pitch of the cones
by RCWA simulation.

Flip-chip type blue LEDs on MPSS, PSS and FSS with anarea of 600 µm x 500 µm were prepared in order to compare their light output. The light output improvement ratiosof the blue LED on the FSS at 50 mA are summarized inFig. 7. The blue LEDs on MPSS with a pitch of 600 nm
had the highest light output power at 50 mA among thetested samples.

4 Conclusions
In this paper, we have presented the advantages of the
MPSS over the existing PSS and FSS technology. The
submicron-scale periodic cones of the MPSS can, according to the RCWA simulation, enhance the transmissivity at
the GaN/sapphire interface thanks to the diffraction effect.
TEM and CL observation results clearly show that the
MPSS can provide a thinner GaN template of equal or
higher quality than PSS. A 3 μm-thick high quality GaN
template on the MPSS with a dislocation density of 1.9 x
108 cm-2 has been demonstrated And the LEDs on MPSS
with a 600 nm pitch show the highest light output power
among the evaluated LEDs on various types of substrates
as it is 1.89 times higher than that on FSS and 1.05 times
higher than that on PSS. The pitch dependence of the light
output improvement is in excellent agreement with that of
the simulated transmissivity at the GaN/sapphire interface.
As a result of the comparisons made of the GaN templates
and LEDs on MPSS, PSS and FSS, we have come to theconclusion that the MPSS does provide the most cost effective solution for high performance LEDs.
Acknowledgements This research has been supported by a
NEDO project.

References

Figure 7 Light output improvement ratios of the LEDs on MPSS
and PSS compared to the LED on FSS

Furthermore, the blue LEDs on MPSS with a pitch of 600
nm had a light output power1.89 times higher than that on
the FSS, and a light output power 1.05 times higher thanthat on the PSS. However, since the GaN templates oneach of the MPSS, PSS and FSS used in the experiments
had nearly the same TDD, all the LEDs should have nearly
the same IQE. Therefore the improvement of the MPSS
should attribute to the enhancement of LEE by the diffraction effect.
The results show that the pitch dependence of the light
output improvement is in good agreement with that of thesimulated transmissivity at the GaN/sapphire interface.
On the other hand, the reason why the calculation values
and the experiment results differ is that the RCWA simulation only shows results of a one-way transmissivity of the© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim[1] T. Seko, S. Mabuchi, F. Teramae, A. Suzuki, Y. Kaneko, R.
Kawai, S. Kamiyama, M. Iwaya, H. Amano, and I. Akasaki,
Proc. SPIE, Photonics West 2009, to be published.
[2] C. C. Bernhard, Endeavour 26, 79 (1967).
[3] H. Toyota, K. Takahara, M. Okano, T. Yotsuya, and H. Kikuta, Jpn. J. Appl. Phys. 40, L747 (2001).
[4] R. Kawai, T. Mori, W. Ochiai, A. Suzuki, M. Iwaya, H.
Amano, S. Kamiyama, and I. Akasaki, Phys. Status Solidi C 6,
S830 (2009).
[5] H. J. Kang, S. U. Cho, E. S. Kim, and M. Y. Jeong, Proc.
SPIE, Photonics West 2013, to be published, February
2013/Vol. 52(2).
[6] A. Ishihara, R. Kawai, T. Kitano, A. Suzuki, T. Kondo, M.
Iwaya, H. Amano, S. Kamiyama, and I. Akasaki, Phys. Status
Solidi C 7, 2056 (2010).
[7] M. G. Moharam, Proc. SPIE 883 (1988).
[8] M. G. Moharam and T. K. Gaylord, J. Opt. Soc. Am. 71,(1981).
[9] H. Yoshida, M. Kuwabar, Y. Yamashita, K. Uchiyama, and H.
Kan, Appl. Phys. Lett. 96, 211122 (2010).

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Applied Physics Express 7, 034104 (2014)
http://dx.doi.org/10.7567/APEX.7.034104

Multijunction GaInN-based solar cells using a tunnel junction
Hironori Kurokawa1, Mitsuru Kaga1, Tomomi Goda1, Motoaki Iwaya1*, Tetsuya Takeuchi1,
Satoshi Kamiyama1, Isamu Akasaki1,2, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
E-mail: iwaya@meijo-u.ac.jp
2

Received January 10, 2014; accepted February 14, 2014; published online March 3, 2014
We fabricated and characterized a two-junction GaInN-based solar cell using a tunnel junction fabricated by crystal growth. This solar cell has two
active layers with a differing bandgap energy corresponding to blue or green light. We conﬁrmed that the open-circuit voltage (VOC) in this solar cell
was increased by the series connection using the tunnel junction. The short-circuit current density, VOC, ﬁll factor, and energy conversion efﬁciencyof this solar cell were 0.28 mA/cm2, 3.0 V, 0.5, and 0.41%, respectively, under an air mass ﬁlter of 1.5 G at 1-sun irradiation and room temperature.
© 2014 The Japan Society of Applied Physics

roup III nitride semiconductors are widely used inlight-emitting diodes (LEDs)1–4) and laser diodes.5,6)
Since the bandgap of GaInN ternary alloys covers
a broad range of values from 0.65 to 3.43 eV,7) these alloys
are suitable for solar-cell applications.8–10) In a multijunction
solar cell, materials with a broad range of bandgap values
offer many possibilities. To date, there have been several
reports of nitride-based solar cells11,12) and we havepreviously succeeded in fabricating GaInN-based solar
cells.13–18) By improving the crystal quality of GaInN
through the application of GaInN superlattice (SL) structures,14,15) the conversion efﬁciency of GaInN-based solar
cells has been increased. Moreover, an energy-conversion
efﬁciency (©) of up to 4.0% for GaInN-based solar cells
under an air mass ﬁlter of 1.5 G and 300 suns irradiation
has been achieved at room temperature (RT).18)
In contrast, the realization of multiple junctions in GaInNbased solar cells remains an important challenge. Tunneljunction formation using crystal growth19,20) and mechanicalstack21) methods has often been used to fabricate multiple
junctions in material systems such as GaAs, Ge, and
AlGaInP. In particular, tunnel-junction fabrication by crystal
growth has achieved a high energy-conversion efﬁciency insolar cells. A multijunction GaInN-based solar cell was previously fabricated by the mechanical stack method.22) On theother hand, there is no report of a multijunction nitride-basedsolar cell using a tunnel junction fabricated by crystal growth.
In this study, we devised a two-junction GaInN-basedsolar cell using a tunnel junction fabricated by crystal growth.
We conﬁrm that the open-circuit voltage (VOC) in this solar
cell was increased by the series connection using the tunnel
junction.
All samples were grown by metalorganic vapor-phase
epitaxy (MOVPE). Figure 1 shows a schematic view ofthe sample structure. Trimethylindium, trimethylaluminum,
trimethylgallium, triethylgallium, and ammonia were used as
the source gases. We also used ethyl-bis(cyclopentadienyl)magnesium (Et-Cp2Mg) and silane (SiH4) as dopant precursors for p- and n-type layers, respectively. C-planesapphire was used as the substrate. The device structure wasdesigned as a GaInN-based solar cell with two GaInN-based
SL active layers on a sapphire substrate covered with a lowtemperature buffer layer (LT-buffer layer) grown at 550 °C.23)
The device structure consisted of the LT-buffer layer, an
unintentionally doped GaN underlying layer (³3 µm), a ﬁrst

G

2nd p-type GaN layer (80 nm)
2-pairs Ga0.7In0.3 N (2.5 nm) /GaN (12 nm) SL
2nd active layer
2nd n-type GaN (200 nm) layern++ GaN (30nm)
p++ Ga 0.8 In 0.2 N (3nm)

Tunnel junction

1st p-type GaN (100nm) layer
p-type Al 0.18 Ga 0.82 N (20 nm) cladding layer
5-pairs Ga0.83 In 0.17 N (3 nm) /GaN (10 nm) SL
1st active layer

1st n-type GaN (2 μm) layer
Unintentionally-doped GaN
LT-buffer layer
C-plane Sapphire Substrate

Fig. 1. Schematic view of the sample structure.

n-type GaN layer (³2 µm) grown at 1,050 °C, a 5-pair
unintentionally doped Ga0.83In0.17N (3 nm)/GaN (10 nm) SL
ﬁrst active layer grown at 780 °C, a p-type Al0.18Ga0.82N
cladding layer (20 nm) grown at 780 °C to suppress thermal
damage at a second n-type GaN layer,24) a ﬁrst p-type GaN
layer (100 nm) grown at 950 °C, the tunnel junction [p++type Ga0.8In0.2N (3 nm) and n++-type GaN (30 nm)] grown
at 780 °C, a second n-type GaN layer (200 nm) grown at
1,000 °C, a 2-pair unintentionally doped Ga0.70In0.30N (2.5nm)/GaN (12 nm) SL second active layer grown at 730 °C,
a second p-type GaN layer (80 nm) grown at 800 °C, and ap+-type GaN contact layer (5 nm) grown at 800 °C. The Si
and Mg concentrations in the n++-type GaN and p++-type
Ga0.8In0.2N layers were set to 4 © 1020 and 1 © 1020 cm¹3,
respectively. The properties of the tunnel junction, such as thedoping concentrations, ﬁlm thickness, and InN molar fraction,
were based on the results of Ref. 25. Further, in order toobtain a high short-circuit current density (JSC) in GaInNbased solar cells, the thickness of the GaInN active layer was

034104-1© 2014 The Japan Society of Applied Physics

Appl. Phys. Express 7, 034104 (2014)

H. Kurokawa et al.

Measurement

X-ray intensity

required to be greater than 150 nm.13) Finally, a wide-gap
material top cell is required to fabricate a high-performance
multijunction solar cell.19,20) However, the growth conditions
to achieve a high crystallinity in a GaInN-based active layerof such thickness are difﬁcult to determine even in singlejunction solar cells. Because the purpose of this study is the
fabrication of a multijunction GaInN-based solar cell, we
designed a structure with relatively few periods in the GaInNbased SL active layer. Similarly, we designed the InN molar
fraction in the ﬁrst SL active layer to be lower than that in thesecond SL active layer for ease of fabrication, because lowtemperature growth is required to realize the high InN molar
fraction GaInN in MOVPE. Moreover, the p+-type GaN
contact layer, with a Mg concentration of approximately
1 © 1020 cm¹3, was introduced to reduce the contact resistance between the p-type electrodes.26)
After annealing the samples at 550 °C for 10 min inambient air to activate the p-type layer, we fabricated a faceup type solar cell. A 100-nm-thick indium tin oxide (ITO)
p-type ohmic transparent electrode and a Ti (20 nm)/Au (150nm) pad electrode were deposited onto the p+-type GaN
contact layer. After annealing the ITO p-type electrodeat 550 °C for 10 min in ambient N2 to reduce the contact
resistance between the p+-type GaN contact layer and the
ITO p-type electrode, a Ti (30 nm)/Al (100 nm)/Ti (20 nm)/Au (150 nm) ohmic n-type electrode was deposited onto theﬁrst n-type GaN layer, which was exposed by inductively
coupled plasma etching. Finally, we performed annealing
at 525 °C for 10 min in ambient N2 to reduce the contact
resistance between the ﬁrst n-type GaN layer and the Ti/Al/
Ti/Au n-type electrode. In this study, the ITO p-type electrode was deposited by sputtering, while the other electrodes
were deposited by electron-beam evaporation; no antireﬂection coating was used for this device. The device andp-type pad electrode dimensions in this two-junction GaInNbased solar cell were 350 © 350 and 100 © 100 µm2,
respectively.
The current density versus voltage characteristics in solarcell mode (J–V curve in solar-cell mode) of this device were
measured using an air mass 1.5 G (AM 1.5 G) solar simulator(Asahi Spectra HAL-320) with a light intensity of 1 sun at
RT. We also characterized the injection-current density versus
voltage relation in LED mode (J–V curve in LED mode) andspectra from this device using a semiconductor parameter
analyzer (Agilent Technologies 4156C) and spectrum analyzer (Hamamatsu Photonics TM-UV/VIS C100820AH). The
concentrations of Mg and Si in the sample were measured bysecondary-ion mass spectrometry. The thickness and InN
molar fraction of each GaInN/GaN SL active layer werecharacterized by X-ray diffraction (XRD; Rigaku SmartLab)using an XRD simulator (Rigaku GlobalFit). The surface
morphology of this device was characterized by atomic
force microscopy (AFM) and scanning electron microscopy(SEM). In addition, interfaces at the tunnel junction wereobserved by cross-sectional transmission electron microscopy(TEM).
Figure 2 shows the XRD (0002) 2ª/½ scan spectra
obtained from this device and the simulation results using
design values. This ﬁgure displays clear satellite peaks fromthe ¹4th to the +2nd order, indicating that the two SL active
layers are relatively well formed even in the multijunction

Simulation
32

33

34
35
2θ [degree]

36

37

Fig. 2. XRD (0002) 2ª/½ scan spectra obtained by measurement andsimulation using the design values.

Measurement area
10 × 10 µm2

RMS= 2.26 nm

5μm
(a)

2μm
(b)
Fig. 3. Plan-view (a) AFM and (b) SEM images of the device.

region. Moreover, the consistency between the experimental
XRD spectra and the simulation results conﬁrm that thisdevice was fabricated as designed. However, there were
problems related to the surface structure. Figure 3 shows theplan-view AFM and SEM images, where a high density(1.6 © 109 cm¹2) of growth pits of is apparent on the device
surface. Moreover, a root-mean-square (RMS) surface roughness of approximately 2.3 nm was measured by AFM.
Considering the RMS surface roughness is usually lower
than 1 nm in a blue LED, the measured RMS value of 2.3 nm
is high. This results from deterioration by the multijunction
formation. From the viewpoint of the surface morphology,
because the largest pit diameter is of the order of several tensof nanometers, these pits are likely to have originated fromthe second p-type GaN layer. Figure 4 shows the crosssectional bright-ﬁeld TEM image of this device, conﬁrming

034104-2© 2014 The Japan Society of Applied Physics

H. Kurokawa et al.

Emission intensity[arb. units]

Appl. Phys. Express 7, 034104 (2014)[0001]pit
2nd p-type GaN layer[11-20]pit[1-100]

2nd SL active layer
2nd n-type GaN
Tunnel junction

1st p-type GaN
p-type AlGaN cladding layer
1st SL active layer
1st n-type GaN layer

200nm

60k

RT, DC
20mA injection

50k
40k

350 μm

30k
20k
10k
0
350

400

450 500 550
Wavelength[nm]

600

650

Fig. 4. Cross sectional bright-ﬁeld TEM image of the device.

Current density [mA/cm2]

Fig. 6. Emission spectrum and emission microscopic image (inset) of thedevice for an injection current of 20 mA.

103
102
10

RT, DC

1

100
10-1
10-2
10-3
10-4
-8

-6

-4

-2 0 2
Voltage [V]

4

6

8

Fig. 5. J–V curve for the device in LED mode.

that the thickness of each layer was as designed; we also
conﬁrmed that each SL structure was well formed. In
contrast, many new dislocations were introduced near thetunnel junction, probably as a result of the high concentrationof impurities in the tunnel junction. Further, growth pits wereformed in the second p-type GaN layer, as observed in the
TEM image (Fig. 4) as well as in the results of plan-view
SEM and AFM (Fig. 3). Therefore, the optimization of thegrowth conditions is essential.
Figures 5 and 6 show, respectively, the J–V curve in LED
mode and the emission spectrum with an emission microscopic image from this device for an injection current of
20 mA (injection current density of ³1.8 © 104 mA/cm2) at
RT. Two emission peaks at 436 and 516 nm, corresponding
to the blue and green wavelengths, respectively, are clearly
apparent in Fig. 6. These emissions presumably correspondto the ﬁrst and second GaInN SL active layers, respectively.
Moreover, the emission wavelengths closely match thewavelengths expected from each GaInN SL active layer.
Therefore, we conﬁrm that the carriers are injected, as
expected, by the tunnel junction. In contrast, the reverse
leakage-current density was lower than 10¹2 mA/cm2 when avoltage of ¹8 V was applied (see Fig. 5). The value of thisreverse leakage current is normally very low in GaInN-basedsolar cells17) and growth pits generally act as leakage-current
sources. However, we believe that the growth pits in thissample did not act as leakage-current sources, given thatthese were generated in the second p-type GaN layer and did
not penetrate the second GaInN SL active layer. In addition,
the forward operated voltage for an injection current densityof 0.3 mA/cm2, a value that almost equals the JSC of thisdevice operated in solar-cell mode with an irradiation of
1 sun, was approximately 3.8 V. Since the turn-on voltage oftypical green and blue LEDs is approximately 2 V, the turnon voltage of this device connected in series is slightly lower
than the theoretical expected value. This decrease in the turnon voltage is correlated to many defects generated by the
fabrication of the tunnel junction. Also, the differential
resistance for an injection current density of 0.3 mA/cm2 in
LED mode was approximately 1.0 k³0cm2. Since the device
resistance for that injection current in blue and green LEDs
is as low as several tens of ohms centimeters squared, thedifferential resistance of this device is extremely high.
Therefore, it is likely that the device resistance was increasedby the high resistance of the tunnel junction. Further, we
conﬁrmed the presence of non-uniform emission-light-dotsfrom Fig. 6. This effect is particularly signiﬁcant in theblue region. We previously reported that hydrogen in p-type
GaN diffuses laterally through the exposed portions of thep-type GaN and not vertically through the n-type GaN layer
above it.27) Therefore, it is suggested that this is caused byinsufﬁcient activation of the p-type Al0.18Ga0.82N cladding
layer, the ﬁrst p-type GaN layer, and the p++Ga0.8In0.2N layer
in the tunnel junction. Further, the device resistance wasincreased as a result of insufﬁcient activation. In addition,
since the growth pits and defects are present on the surfaceof this sample, it is considered that a non-uniform activationof the p-type AlGaN cladding layer, the ﬁrst p-type GaN
layer, and the p++Ga0.8In0.2N layer is a consequence of these
defects. For these reasons, the reduction of the resistancein the tunnel junction optimized by growth conditions andprocessing methods is essential.
Figure 7 shows the J–V curves in solar-cell mode for thistwo-junction GaInN-based solar cell produced by the solar
simulator (AM 1.5 G, 1 sun) at RT. Judging from this ﬁgure,the solar cell exhibited good characteristics. The ©, VOC, ﬁll
factor, JSC, shunt resistance (RSH), and series resistance (RS)of this solar cell were 0.41%, 3.0 V, 0.5, 0.28 mA/cm2,
2.8 k³0cm2, and 0.830 k³0cm2 at 1-sun irradiation, respectively. From these results, the JSC is lower than that of acommon GaInN-based solar cell owing to a thin-ﬁlm active
layer in this solar cell. According to a simple calculation, thetheoretical JSC for a 1-sun irradiation of the active layer with
an absorption edge of 516 or 436 nm is approximately 9.2 or

034104-3© 2014 The Japan Society of Applied Physics

Current density [mA/cm2]

Appl. Phys. Express 7, 034104 (2014)

H. Kurokawa et al.

0.3

Acknowledgments This work was supported by a Grant-in-Aid for
Scientiﬁc Research (C) No. 23560015, Grant-in-Aid for Young Scientists (A)No. 24686003, the New Energy and Industrial Technology Development
Organization (NEDO) Project “Research and Development on Innovative Solar
Cells”, and the Ministry of Education, Culture, Sports, Science and Technology
Supported Program for the Strategic Research Foundation at Private Universities,
Japan, 2012–2016.

0.2

0.1

0.0
0

RT
A.M. 1.5G
1 sun
1

2
Voltage [V]3Fig. 7. J–V curves in solar-cell mode of the two-junction GaInN-basedsolar cell, produced by the solar simulator (AM 1.5 G, 1 sun) at RT.

4.1 mA/cm2, respectively. However, the active layer of thisdevice is not thick enough to absorb light. Because thethicknesses of the GaInN layers are 5 and 15 nm, respectively, the absorbance of each GaInN layer was only 5 and15%, respectively, assuming the absorption coefﬁcient of
GaInN to be 105 cm¹1. Therefore, the theoretical JSC foreach GaInN active layer under a 1-sun irradiation was 0.45
and 0.61 mA/cm2, respectively. Thus, when considering thereﬂection of the semiconductor surface, it is reasonable thatthe resulting JSC is close to the theoretical values. In addition,a relatively high RSH is obtained because the leakage-current
density in this device is small. However, because this value
is insufﬁcient, the reduction of pit and defect densities isrequired for the improvement of GaInN solar cells. In
contrast, RS is approximately one order of magnitude higher
than that in the other GaInN-based solar cells. In addition,
this RS was almost identical to the value of the differential
resistance evaluated in LED mode. Therefore, we concluded
that RS had increased as a result of the high tunnel-junction
resistance and insufﬁcient activation in the p-type layer.
Moreover, we conﬁrmed the presence of two inﬂection points
in the J–V curve. We think that these inﬂection points are dueto the inactivation of carriers that produces a high resistancein the tunnel junction. Further, because RS is high, the ﬁllfactor of the device is low and optimization of the process
and crystal growth conditions to reduce RS is required. The
VOC of GaInN-based single-junction solar cells is approximately 2 V. Therefore, the VOC for the two-junction solar cell
was increased by a factor of approximately 1.5 through theseries connection using the tunnel junction. We previously
reported a strong correlation between the pit density and
VOC.17) We also reported that a pit density below 107 cm¹2 isrequired to obtain a high VOC.17) Therefore, the reduction of
the pit density is essential for further increases in VOC.
In conclusion, we designed and characterized a two-junction GaInN-based solar cell using a tunnel junction fabricatedby crystal growth. We conﬁrmed that the VOC in this solar cell
was increased by the series connection using the tunnel
junction. Although a high resistance in the tunnel junction
remains a problem, we demonstrated that multijunction
GaInN-based solar cells are possible by tunnel-junction
fabrication.

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Lateral Hydrogen Diffusion at p-GaN Layers in Nitride-Based Light Emitting Diodes with
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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JK12
http://dx.doi.org/10.7567/JJAP.52.08JK12Lateral Hydrogen Diffusion at p-GaN Layers in Nitride-Based Light Emitting Diodes
with Tunnel Junctions
Yuka Kuwano1 , Mitsuru Kaga1 , Takatoshi Morita1 , Kouji Yamashita1 , Kouta Yagi1 ,Motoaki Iwaya1 , Tetsuya Takeuchi1 , Satoshi Kamiyama1 , and Isamu Akasaki1;2
1
2Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

Received October 13, 2012; accepted April 24, 2013; published online August 20, 2013
We demonstrated lateral Mg activation along p-GaN layers underneath n-GaN surface layers in nitride-based light emitting diodes (LEDs) withGaInN tunnel junctions. A high temperature thermal annealing was effective for the lateral Mg activation when the p-GaN layers were partly
exposed to an oxygen ambient as etched sidewalls. The activated regions gradually extended from the etched sidewalls to the centers with anincrease of annealing time, observed as emission regions with current injection. These results suggest that hydrogen diffuses not verticallythorough the above n-GaN but laterally through the exposed portions of the p-GaN. The lowest voltage drop at the GaInN tunnel junction was
estimated to be 0.9 V at 50 mA with the optimized annealing condition. # 2013 The Japan Society of Applied Physics

1. Introduction

Nitride semiconductor materials are most promising materials for blue light-emitting diodes (LEDs), ultraviolet laser
diodes, and high power electronic devices.1,2) It had beenvery diﬃcult to obtain p-type GaN until Amano et al.
achieved p-type conduction in Mg-doped GaN treated with
low-energy electron beam irradiation (LEEBI).3–5) Since
then, thermal annealing without hydrogen or ammonia
ambient was also proposed to obtain p-type conduction in
the Mg-doped GaN.6–8) Several groups have reported thatMg acceptors are passivated with hydrogen atoms, resultingin Mg–H bond formation.9–13) During LEEBI and thermal
annealing, the Mg–H bonds are broken and the Mg acceptors
are electrically activated.
Through the above studies, it have been also revealed
that hydrogen atoms exist in Mg-doped GaN, but not in
undoped or Si-doped GaN even after thermal annealing
with hydrogen ambient.14) This phenomena has been also
observed in other semiconductor materials, such as Si and
InP. Furthermore,15,16) it has been suggested that hydrogenatoms can migrate only in p-type materials, but not in n-typematerials.
Meanwhile, a novel current injection in GaInN-based
LEDs using tunnel junctions was proposed in order toavoid the poor conductivity of the p-type GaN.17,18) Aftera growth of a conventional GaInN-based LED, a pþþ Ga(In)N/nþþ -GaN tunnel junction and an n-GaN weresuccessively grown on the LED structure. This structure isuseful for higher light extraction in LEDs and current
conﬁnement in vertical cavity surface emitting lasers due
to excellent current spreading in the top n-GaN layers.
Interestingly, in this structure, p-type layers are fully
covered with n-type layers, and hydrogen atoms in thep-type layers may not pass through the n-type layers fromthe p-type layers.
In this paper, we optimize thermal annealing conditionsfor Mg activation in such a LED structure with the tunnel
junction and the n-GaN surface layer. We then demonstrate
lateral Mg activation from etched sidewalls in the LED
structure, indicating that hydrogen atom can only diﬀuse
laterally in the p-GaN layers due to the n-GaN surface
layers.

p-typelayer

n+-GaN:10nm

n+-GaN:10nm

n-GaN : 500nm
n++-GaN:30nm
p++-GaInN:5nm
p-GaN:100nm
p-AlGaN:25nm

n-GaN : 500nm
n++-GaN:30nm
p++-GaInN:5nm
p-GaN:100nm
p-AlGaN:25nm

TunnelJunction

GaInN/GaN 5QW:75nm

GaInN/GaN 5QW:75nm

n-GaN:2.4µm

n-GaN:2.4µm

u-GaN : 3µm

u-GaN : 3µm

LT-GaN

LT-GaNC-plane Sapphire Sub

C-plane Sapphire Sub

(a)

(b)

(c)
Fig. 1. (Color online) Schematics of (a) the wafer structure containing
the tunnel junction and the LED structure and (b) the same wafer structure
but after the etching. The arrows indicate potential hydrogen diﬀusion paths
during the Mg activation.

2. Experiments

We prepared the wafer containing a tunnel junction structure
and a LED structure in series as shown in Fig. 1(a). After
deposited a 25-nm-thick GaN layer at low temperature ona sapphire substrate, a 3.0-m-thick undoped GaN, a 2.4m-thick Si-doped GaN bottom contact layer, a multiplequantum well active layer consisting of ﬁve periods of
3 nm GaInN/12 nm GaN (emission wavelength: 450 nm),a 25-nm-thick Mg-doped AlGaN, and a 100-nm-thick Mgdoped GaN were grown by metalorganic vapor phase
epitaxy. Up to the Mg-doped GaN layer, the structure is
the same as a conventional LED structure. Additionally, atunnel junction, a 500-nm-thick Si-doped GaN for lateralcurrent spreading, and a 10-nm-thick highly Si-doped GaN

08JK12-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JK12

Y. Kuwano et al.

Fig. 3. (Color online) Emission regions of the LEDs with the tunnel
junctions activated under the various thermal annealing conditions. The caseof a standard LED without the tunnel junction but with a semi-transparent
Ni/Au electrode annealed under the standard condition is also shown.
Fig. 2. (Color online) Mg, Si, and In depth proﬁles at the tunnel junctions
measured by SIMS.

top contact layer were successively grown on the structure.
The tunnel junction consists of a 5-nm-thick heavily Mgdoped Ga0:8 In0:2 N layer and a 30-nm-thick heavily Si-doped
GaN layer. Triethylgallium (TEG), ammonia (NH3 ), ethylbis(cyclopentadienyl)magnesium (Et-Cp2 Mg), and silane(SiH4 ) were used as Ga, N, Mg, and Si sources, respectively.
Figure 2 shows depth proﬁles of Mg, Si, and In atoms in
the wafer by Secondary Ion-microprobe Mass Spectrometry.
1  1020 cm3 Mg and 3  1020 cm3 Si were observed at
the tunnel junction. A band diagram of structure of Fig. 1(a)shown in Fig. 1(c).
Then LEDs with tunnel junctions were fabricated from thewafer mentioned above by a newly modiﬁed process basedon our standard LED process. In the modiﬁed process, two
diﬀerent parts exist comparing with the standard LED
process. First, the thermal annealing was carried out afteretching to expose the Si-doped GaN bottom layer surface
for the bottom n-contact. Typically we perform the thermal
annealing before the etching in a standard LED process. The
reason for changing the order of these process steps is to
expose a portion of the p-layers as an etched sidewall andform a potential hydrogen diﬀusion path during the Mg
activation as shown in Fig. 1(b) by arrows. For the thermal
annealing conditions, several annealing temperatures (525,625, 725  C) and holding times (4, 15, 30 min) were selected.
Our typical annealing conditions in the standard process are
525  C and 4 min. Oxygen gas was used as an ambient gas.
Second, a single patterning process for forming metalelectrodes was carried out since both the electrodes weren-contacts in the LEDs with tunnel junctions. Furthermore,no transparent electrode is required since the Si-doped GaN
top layer act as current spreading layer, like a semitransparent electrode. A size of the LED was 300  300 m2 .
A standard LED (without tunnel junction) with a semitransparent Ni/Au electrode was also fabricated for comparison. Emission regions and current–voltage characteristics of
the LEDs were measured under direct current conditions atroom temperature. Voltage drops at the tunnel junction
regions were estimated by subtracting voltage drops of thestandard LED from those of the LEDs with tunnel junctions.
3. Results and Discussion

Figure 3 shows the emission regions of LEDs with thetunnel junctions activated under the various thermal
annealing conditions. While a uniform emission region
was obtained from the LED with the tunnel junction
annealed at 725  C for 30 min, the emission regions of theother LEDs with the tunnel junctions were restricted near
the etched sidewall of the LEDs. More precisely narrower
emission regions just adjacent to the etched sidewall of
the LEDs were obtained in the cases of lower annealing
temperature and shorter annealing time. Then with increasesof annealing temperature and annealing time, the emission
regions progressively extended to the center of the LEDs.
We conﬁrmed that the standard LED with a semi-transparent
Ni/Au electrode showed uniform emission regions, so theactive regions in all the LEDs should be uniform. Thus, thenonuniform emission should be caused by nonuniformity ofcurrent injections at the tunnel junctions. We then supposethat Mg acceptors are activated ﬁrst at the etched sidewallsof the LEDs where Mg-doped layers are directly exposed tothe oxygen ambient, and then Mg activation regions extend
from the edge to the center.
To discuss the above point further, Mg activation length
was plotted as a function of the square root of annealing
time as shown in Fig. 4. Here Mg activation length isdeﬁned as shown in Fig. 4(a), corresponding to the length ofthe emission region. As shown in Fig. 4(a), we measured
total six points which were away from the corners and the
top n-electrode in order to minimize inﬂuences of thecorners and the electrode. In the Fig. 4(b), the dots indicate
the average values of the six points and the error bars
indicate the range of the values. As shown in Fig. 4(b), theMg activation length is proportional to the square root of
annealing time. Therefore, the Mg activation length should
be governed by some sort of diﬀusion process. The diﬀusion
coeﬃcient D was estimated to be 9:6  109 cm2 /s from

08JK12-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JK12

Y. Kuwano et al.

725  C, the voltage drops at the tunnel junctions was clearlydecreased. At the same time, no annealing time dependenceof the voltage drop was observed within the range of theexperiments. The lowest voltage drop at the GaInN tunnel
junction is 0.9 V at 50 mA in the experiments.
4. Conclusions

(b)

(a)

Fig. 4. (Color online) (a) Mg activation length deﬁned as shown witharrows. (b) Mg activation length as a function of square root of annealing
time at 725  C.

8annealing temperature: 725[°C]7Voltage [V]

6
5
4
3

4min
15min
30min
standard LED w/o tunnel junction

2
1
0

5

10

15

20

25

30

35

40

45

50

Acknowledgements

Current [mA]

This work was supported by Grant-in-Aid for ScientiﬁcResearch (C) (No. 23560015), and the Ministry of Education, Culture, Sports, Science and Technology Supported
Program for the Strategic Research Foundation at Private
Universities, 2012–2016. The SIMS measurement was
carried out by courtesy of Toray Research Center, Inc.

(a)8annealing time: 30min7
6

Voltage [V]In summary, we have optimized thermal annealing conditions for the Mg activation in the LED structure with thetunnel junction and the n-GaN surface layer. We found thatthe p-GaN layers underneath the n-GaN surface layers wereactivated when the p-GaN layers were exposed to an oxygenambient as etched sidewalls. Higher temperature and longer
time for the thermal annealing resulted in an extension of
Mg activated regions from the etched sidewall to the center.
This result suggest that hydrogen diﬀused not through theabove n-GaN layers but along p-GaN layers toward theetched sidewalls. We also found the annealing temperatureaﬀects a degree of the Mg activation and the annealing timeaﬀects a length of the Mg activation from the exposed
portion. The lowest voltage drop at the GaInN tunnel
junction was estimated to be 0.9 V at 50 mA with the
optimized annealing condition.

543525[°C]625[°C]725[°C]
standard LED w/o tunnel junction

2
1
0

5

10

15

20

25

30

35

40

45

1) H. Amano, N. Sawaki, I. Akasaki, and Y. Toyoda: Appl. Phys. Lett. 48(1986) 353.
2) S. Nakamura, T. Mukai, and M. Senoh: Appl. Phys. Lett. 64 (1994) 1687.
3) H. Amano, M. Kito, K. Hiramatsu, and I. Akasaki: Jpn. J. Appl. Phys. 28

50

Current [mA]

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(b)

Lett. 81 (2002) 3747.

Fig. 5. (Color online) Current–voltage characteristics of the LEDs withthe tunnel junction annealed under (a) various annealing times at 725  C and(b) various annealing temperatures for 30 min. The case of a standard LEDwith a semi-transparent Ni/Au electrode is also plotted.

5) Y. Okamoto, M. Saito, and A. Oshiyama: Jpn. J. Appl. Phys. 35 (1996)

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Fig. 4(b). At this moment we conclude that hydrogen atoms
diﬀused only toward lateral direction through the etchedsidewalls, but not vertically to the surface since the n-GaN
surface layers blocked the hydrogen diﬀusion. The phenomena that n-layers block the hydrogen diﬀusion is alreadyreported in the other semiconductor materials.9–11)
Figure 5 shows current–voltage characteristics of theLEDs with the tunnel junction under various thermal
annealing conditions. Figure 5(a) shows an annealing timedependence under 725  C and Fig. 5(b) shows an annealing
temperature dependence for 30 min. The result of thestandard LED with a semi-transparent Ni/Au electrode isalso plotted for comparison. The voltage diﬀerence between
the standard LED and the LED with tunnel junction
corresponds to the voltage drop at the tunnel junction.
When the annealing temperature was increased from 525 toAppl. Phys. Lett. 67 (1995) 2666.
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Appl. Phys. Lett. 69 (1996) 3725.

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12) B. A. Hull, S. E. Mohney, H. S. Venugopalan, and J. C. Ramer: Appl. Phys.

Lett. 76 (2000) 2271.
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S. Sakai: Jpn. J. Appl. Phys. 38 (1999) 631.
14) A. Y. Polyakov, N. B. Smirnov, S. J. Pearton, F. Ren, B. Theys, F. Jomard,

15)
16)
17)
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08JK12-3

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Nikitina: Appl. Phys. Lett. 79 (2001) 1834.
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748.
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Chew, and D. E. Sykes: Appl. Phys. Lett. 53 (1988) 758.
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Rajan: Appl. Phys. Lett. 97 (2010) 203502.
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Appl. Phys. Lett. 79 (2001) 2532.

# 2013 The Japan Society of Applied Physics

Dislocation density dependenceof stimulated emission characteristics
in AlGaN/Al multiquantum wells

solidistatuspss

physicaPhys. Status Solidi C 10, No. 11, 1537– 1540 (2013) / DOI 10.1002/pssc.201300265

cwww.pss-c.com

current topics in solid state physics

Yuko Matsubara*,1, Tomoaki Yamada1, Kenichiro Takeda1, Motoaki Iwaya1, Tetsuya Takeuchi1,
Satoshi Kamiyama1, Isamu Akasaki1,2, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8062, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8062, Japan
2

Received 23 June 2013, revised 18 September 2013, accepted 25 September 2013
Published online 28 October 2013
Keywords stimulated emission characteristics, deep UV, AlGaN/AlN, dislocation density* Corresponding author: e-mail 123434041@ccalumni.meijo-u.ac.jp, Phone: +82 52 832 1151, Fax: +81 52 832 1244

We investigated dislocation density dependence on stimulated emission characteristics of AlGaN/AlN multiquantum wells (MQWs: emission wavelength of approximately 285 nm) on AlN templates with various dislocation densities. We found that the stimulated emission

characteristics of the deep UV MQWs were strongly dependent on the dislocation densities. A reduction of thedislocation densities is very important in order to realize
good stimulated emission characteristics of deep UV
MQWs.

© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction AlGaN ternary nitride semiconductors have much potential for optoelectronic applications,particularly in the deep UV/UV region. Recently numerous
reports on the deep UV/UV light emitting diodes havebeen published [1-5]. Especially the UV lasers are important in terms of many applications in the medical, chemical,
and engineering fields. However, developments of the UV
laser diodes [6-13] are hampered by poor material qualityof AlGaN-based heterostructures and difficulty of current
injection, especially with high AlN molar fraction. Opticalpumping is a flexible and convenient excitation techniquefor investigation of stimulated emission characteristics
compared to electrical injection. Most studies on the stimulated emission characteristics in AlGaN by optical pumping have been focused on the optimization of growth conditions and structure designs [14, 15], whereas the stimulated emission characteristics of the AlGaN-based active
layer in conjunction with dislocation density have not been
investigated in detail.
In this study, we investigated the dislocation density
dependence on the stimulated emission characteristics
of AlGaN/AlN multiquantum wells (MQWs) on a thick
AlN templates with various dislocation densities by excita-tion density dependent photoluminescence (PL) measurement.
2 Experiments and results
2.1 MQWs on AlN All samples were grown by hightemperature metalorganic vapor phase epitaxy [16] on a cplane sapphire substrate. The epitaxial wafers were thenpolished and cleaved in the direction of the a-axis for AlN
with a 500 μm cavity length. Figure 1 shows a schematic
view of the sample structure. We prepared the samples
with an emission wavelength of approximately 280 nm.
The quantum well structures are 10-pairs of Al0.25Ga0.75N(4 nm)/AlN (8 nm) MQWs. The sapphire was polished toless than 100 μm in thickness, and then cleavages weremade using a diamond scriber. The samples have cleaved
mirror facets without coating.

© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss

1538

c

Y. Matsubara et al.: Dislocation density dependence of stimulated emission characteristics in AlGaN/AlN MQWs

We determined the layer thickness and the AlN molar Table 1 Dislocation density of AlN templates
fraction of each AlGaN MQWs from the satellite peaks ofthe 2θ/ω scan X-ray diffraction profiles. The threading disSample A
Sample B
Sample C
location densities of AlN underlying layers were controlled
Edge disloca910-2from 3×10 to 1×10 cm by changing the growth condi5.51×109
1.00×1010
2.95×109
tion density
tion. The value of the dislocation density of each sample[cm-2]was estimated from the values of full width half maximum
Mixed disloca(FWHMs) of the X-ray rocking curve (XRC) ω-scan
2.31×108
5.29×107
7.85×107
tion density(0002) and (10-12) diffractions with the following formula[cm-2][17]:tion density values correlated with optical properties. Notethat a cross-check was conducted using plan-view transb2
N =
(1)
G2 ,
mission electron microscopy and cathodoluminescence re4 .3 5 bgarding the dislocation density.
where N is the threading dislocation density, b is the magnitude of Burgers vector, and β is the half width by measuring the omega magnitude of XRC. Table 1 shows thedislocation density of each AlN template. We derived thedensity values of edge dislocations and mixed dislocations,which should act as non-radiative recombination centers[18]. Therefore, we here used the sum of both as disloca-

2.2 Measurement We characterized the stimulated
emission of the MQWs by excitation density dependent
photoluminescence measurement at room temperature(RT). The samples were excited by the 4th harmonic (266nm) of the Q-switched YAG:Nd laser radiation. The pulse
width was 5 ns and the frequency was 10 Hz. A neutral
density filter was used to control the excitation density.
The stimulated emission measurements were performed inedge-detection mode. The light beam from the YAG laser
was refracted by a prism and irradiated on the samples.
The photoluminescence light emitting from the sample
edge was passing through a receiving lens, and detected in
the spectrometer (SPEC Miniature Fiber Optic Spectrometer). We focused the YAG laser on the sample surface. The
irradiation area on the surface was about 0.015 cm2, andthe shape was circle. Note that in these samples only thewell layers were selectively excited.

Figure 1 Schematic view of sample structure.

289nm
25k

2.2

RT

20k

1.5

285

30k

2.9

20k

15k

2.0

RT

10k

1.4

289nm

2.6

25k

20k

2.2

RT

15k

10k

1.7

5k

0.49
0
280

2
MW/cm

285nm
Emission intensity [arb. units]

30k

10k

2
MW/cm
Emission intensity [arb. units]

Emission intensity [arb. units]

40k

2
MW/cm
2.7

5k

0.49
290

295

300

Wavelength [nm]

0
270

280

290

300

Wavelength [nm]

310

0
270

0.49
280

290

300

Wavelength [nm]

310

(a)
(b)
(c)
Figure 2 PL spectra under different excitation power densities. (a) Dislocation density of 3×109 cm-2. (b) Dislocation density of
5×109 cm-2. (c) Dislocation density of 1×1010 cm-2.
© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

www.pss-c.com

Contributed
ArticleDislocation density
3x109 cm-2

30
FWHM [nm]

1539

25

5x109 cm-2
1x1010 cm-2

20
15
10
5
0
0.0

0.5

1.0

1.5

2.0

2.5

2
Excitation power density [MW/cm ]Emission intensity [arb. units]

Phys. Status Solidi C 10, No. 11 (2013)

3E9

80k

RT

70k

5E9

60k
50k

1E10

40k

Dislocation density
9 -2
3x10 cm
9 -2
5x10 cm
10 -2
1x10 cm

30k
20k
10k
0

0

1
2
3
4
5
Excitation power density [MW/cm2]

Figure 3 FWHM of the PL emission intensity as a function of
the excitation power density.

Figure 4 PL emission peak intensity as a function of the excitation power density.

Figures 2(a)-(c) show the measured emission spectra
under different excitation power densities, varying from 49
kW/cm2 to about 3.0 MW/cm2 at RT for each MQW. In thecase of sample A, the FWHM was approximately 1.5 nm at
pumping power of 2.7 MW/cm2. This FWHM value is limited by a resolution in our experiment equipment. Thus, weconsidered that the stimulated emission was observed fromsample A even though the FWHM was not very narrow.
Figure 3 summarizes the FWHMs of the emission
spectra as a function of the excitation power density. The
FWHM of sample A reduced largely when excitation power density exceeded the threshold power density. From
this figure, the FWHMs of sample B and C are similarly
decreased to the FWHM of sample A. Thus, we concluded
that these samples have reached the stimulated emission.
Figure 4 summarizes the PL emission peak intensity as
a function of the excitation power density. For all samples,
new narrow peaks appear in the spectra as the excitation
power densities are increased above certain values. With
the further increase in the excitation intensity, the intensityof the new peak increases superlinearly. We observed thatthe stimulated emission characteristics of the deep UV
MQWs were strongly dependent on the dislocation densities. As the dislocation density decreased in AlN template,
the threshold power density also decreased. The threshold
power density of the stimulated emission from each AlGaN MQW changes from approximately 500 kW/cm2 to 5
MW/cm2 when the dislocation density changes from 3×109
to 1×1010 cm-2. Thus, a reduction of the dislocation densities is very important in order to realize superior stimulated
emission characteristics of the deep UV MQWs.

der to realize a low-threshold laser diode, a base layer withlow dislocation density is required.

3 Summary We analysed the stimulated emission
characteristics of the AlGaN/AlN MQWs on the AlN underlying layer with the various dislocation densities by excitation density dependent PL measurement. The threshold
power density of each AlGaN MQW changed when thedislocation density changed. Thus, it was found that, in or-

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Acknowledgements This study was supported by the
MEXT-Supported Program for the Strategic Research Foundation
at Private Universities, 2012-2016, and the Adaptable and Seamless Technology Transfer Program through Target-driven R&D,
JST (AS2421230J).

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Schowalter, Appl. Phys. Lett. 91, 051116 (2007).
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Russel, and D. Slater, J. Cryst. Growth 272, 242 (2004).
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[11] H. Tsuzuki, F. Mori, K. Takeda, M. Iwaya, S. Kamiyama, H.
Amano, I. Akasaki, H. Yoshida, M. Kuwabara, Y. Yamashita, and H. Kan, J. Cryst. Growth 311, 2860 (2009).
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Carrier Injections in Nitride-Based Light Emitting Diodes Including Two Active Regions with
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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JG02
http://dx.doi.org/10.7567/JJAP.52.08JG02

Carrier Injections in Nitride-Based Light Emitting Diodes Including Two Active Regionswith Mg-Doped Intermediate Layers
Kenjo Matsui1 , Koji Yamashita1 , Mitsuru Kaga1 , Takatoshi Morita1 , Tomoyuki Suzuki1 ,Tetsuya Takeuch1 , Satoshi Kamiyama1 , Motoaki Iwaya1 , and Isamu Akasaki1;2
1
2Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

Received October 13, 2012; revised December 6, 2012; accepted December 7, 2012; published online May 20, 2013
We achieved simultaneous two emissions from a single light emitting diode containing violet and blue GaInN/GaN quantum well active regions
separated with an intermediate layer. By adjusting a thickness of a Mg-doped region in the intermediate layer, an intensity ratio between violet andblue emissions was changed, caused by different carrier distributions in the two active regions. The intensity ratio from two active regions was alsochanged by changing an amount of current injection. An unintentional Mg incorporation into the active region above the Mg-doped intermediate
layer was observed, which is due to Mg memory effect. # 2013 The Japan Society of Applied Physics

p-electrode

1. Introduction

Optoelectronic devices handling with broad spectra, likesolar light, have been demanded, such as solid state lighting
and solar cells. GaN-based material system is most promising to cover the solar spectrum by using a mixed crystal of
AlN, GaN, and InN.1) However semiconductor optoelectronic devices, such as light emitting diodes (LEDs), laser
diodes, and photodiodes fundamentally deal with just narrowspectra corresponding to bandgaps at active regions. Thus apractical white LED has been fabricated by a combination ofa blue LED and a yellow phosphor material.2)
Recently multi-color LEDs with multiple peak wavelengths have been demonstrated on target to a monolithic
white LED. Various approaches and structures for themultiple peak emissions have been proposed, such as atandem structure with a tunnel junction,3,4) an active region
consisting of various quantum wells (QWs),5–11) and a QW
on various micro facets.12,13) In any case, these devices havea single active region in a single pn junction in principle, likea conventional LED.
Here we are interested in multiple active regions in asingle pn junction. The multiple active regions consist of two
or more active regions separated with intermediate layers
which thickness is comparable to that of the active region.
This structure is useful not only for multi-color LEDs but
also other light emitting devices. For instance, the structure
has been used as a periodic gain structure (PGS) in a verticalcavity surface emitting laser (VCSEL) to increase an opticalconﬁnement factor.14) Also, the structure should be useful tosuppress the eﬃciency droop issue typically observed in
GaN-based LEDs by diluting carrier concentrations with the
thick active regions.5) Note that in the above examples theactive regions have the same peak wavelength. So far the
structure has been applied to GaAs-based and InP-based
device structures mostly as the PGSs in the VCSELs.15)
There are almost no attempts to investigate the two active
region structures with GaN-based materials. This seems
reasonable in a way since holes in GaN-based materials havevery heavy mass and short diﬀusion length which lead topoor carrier injection into the thick active region.
In this paper, we demonstrate simultaneous two emissions
by current injection from two GaInN QW active regions
separated with a Mg-doped intermediate layer. We foundp-GaN
n-electrodep-side blue activeu-GaN
Mg-doped layer
u-GaN
n-side violet active

intermediate layer

n-GaN
Sapphire Sub.

Fig. 1. (Color online) Schematic view of the two wavelength violet/blue
LED, containing the two active regions separated with the intermediate
layer.

that the Mg doping level in the intermediate layer is a key toobtain the simultaneous emissions. Furthermore current
injection dependences on emission intensities and spectrum
patterns are discussed.
2. Experiments

All the samples used in this study are LED structures grownon c-plane sapphire substrates by metalorganic vapor phase
epitaxy. The sample structure is shown in Fig. 1. Before
the LED structure growth, a 25-nm-thick low-temperaturedeposited buﬀer layer and a 3-m-thick undoped GaN
template were grown ﬁrst. Following the growth of the
GaN template, a 2.5-m-thick n-GaN layer, a three-periodof undoped Ga0:88 In0:12 N/GaN QWs as an ‘‘n-side’’ violet
active region, a 45-nm-thick GaN/p-GaN/GaN intermediate
layer, a three-period of undoped Ga0:8 In0:2 N/GaN QWs as a‘‘p-side’’ blue active region, a 20-nm-thick p-type AlGaN
electron blocking layer, a 80-nm-thick p-GaN layer, and a
10-nm pþ -GaN contact layer were grown. Each GaInN/GaN
3QW active region is composed by 2.5-nm-thick GaInN
wells and 13-nm-thick GaN barriers. Regarding the GaN/
p-GaN/GaN intermediate layer, an intentionally Mg-doped
region is located at the center of the intermediate layer. In
this study, the thickness of the Mg-doped region in theintermediate layer was changed while supplying the same

08JG02-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JG02

K. Matsui et al.

(a)

(b)

1E+22

Mg
In

1E+21

1E+09

1E+22

1E+08

1E+21

Intermediate layerp-side activen-side active

1E+09
1E+08

Intermediate layer

1E+07

p-side active

1E+20

n-side active

1E+07

1E+05

1E+18

1E+04
1E+03

1E+17

1E+06
1E+19

1E+05

1E+18

1E+04
1E+03

1E+17

1E+02
1E+16

1E+02
1E+16

1E+01

1E+15

1E+00
0

50

100 150 200
Depth[nm]In intensity [counts]

1E+19

Mg concentration [cm-3]

1E+20
1E+06
In intensity [counts]

Mg concentration [cm-3]

Mg
In

1E+01

1E+15

250

1E+00
0

50

100 150 200
Depth[nm]250Fig. 2. (Color online) Mg and In depth proﬁles measured by SIMS for the samples (a) with the 14 nm Mg-doped layer and (b) with the 1.6 nm
Mg-doped layer.

amount of bis(ethylcyclopentadienyl) magnesium (EtCp2 Mg) ﬂow in order to change a total amount of Mg in
the intermediate layer. Five diﬀerent samples were preparedin which the Mg-doped layer thickness was changed, 14, 4,
1.6, 0.8, and 0 (corresponding to undoped) nm. At the sametime, the total thickness of the intermediate layers was kept
as 45 nm. After the epitaxial growth, all the samples were
fabricated by our standard LED process. A size of LEDs was
300  300 m2 . A semi-transparent p-electrode at the top of
the device was formed by Ni/Au, while Ti/Al/Ti/Au wasused as an n-electrode.
Mg and In atom depth proﬁles of the samples were
measured by secondary ion mass spectrometry (SIMS).
Room temperature photoluminescence (PL) spectra from the
LED structures were measured to conﬁrm emission wavelengths and intensities of the active regions, which wereindependent on a carrier distribution imbalance caused by acurrent injection. Then emission spectra with several current
injection levels were measured from the same LED
structures at room temperature. In addition current–light
output power (I–L) characteristics were measured under
direct current condition at room temperature.
3. Results and Discussion

Mg and In atom depth proﬁles of the samples with the 14
and 1.6 nm Mg-doped layers were shown in Figs. 2(a) and2(b), respectively. Even though the Et-Cp2 Mg ﬂows were
ﬁxed among the two samples, the Mg concentrations at
the intermediate layers were diﬀerent in the two samples,1  1019 cm3 in the sample with the 14 nm Mg-doped layerand 5  1018 cm3 in the sample with the 1.6 nm Mg-doped

layer. This could be attributed to Mg turn-on delay due tothe Mg memory eﬀect. In addition, a considerable amountof Mg atoms in the range of more than 1  1018 cm3 wereincorporated into the p-side active regions. This Mg turn-oﬀ
delay also came from the Mg memory eﬀect. Note that Mg
incorporation in the GaInN well looks higher than that of
GaN barrier, but this behavior may be caused by a sensitivity
diﬀerence of Mg atoms in GaN and GaInN. At this moment
it is hard to estimate the diﬀerence in the Mg incorporation
between GaInN and GaN.
Room temperature PL spectra of all the samples
consisting of the LED structures are shown in Fig. 3. An
excitation source was a 325 nm He–Cd laser. As shown in
Fig. 3, all the samples basically showed similar spectrum
patterns with intensities in the twofold range. Weak peaks
indicated by solid arrows around 400 to 430 nm from then-side active regions and strong peaks indicated by dashedarrows around 450 to 470 nm from the p-side active regions
were consistently observed. Note that the sample-to-sample
variations of the PL peak wavelength correspond to run-torun variations of growth temperatures at the active regions.
Since the similar PL intensities were obtained from all thesamples, the optical qualities of the active regions among all
the samples could be similar. The reason for much weaker
PL intensity from the n-side violet active region compared tothat from the p-side blue active region is a lower excitation
intensity due to a deeper position of the n-side violet active
region from the surface.
Next, room temperature emission spectra by a current
injection of 20 mA were plotted in Fig. 4. Interestingly, thespectra are very diﬀerent among the samples, and also

08JG02-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JG02

Light output power[a.u.]

K. Matsui et al.

0.5
0.4
0.3
0.2

14 nm
4 nm
1.6 nm
0.8 nm
0 nm

0 nm
14 nm

0.8 nm
1.6 nm
4 nm

0.1
0.0
0 10 20 30 40 50 60 70 80 90 100
Current[mA]

Fig. 5. (Color online) Room temperature I–L characteristics for all thesamples.

Fig. 3. (Color online) Room temperature PL spectra of all the samples
with the LED structures. PL peak positions from the n- and p-side active
regions are indicated by the solid and dashed arrows, respectively.

Fig. 4. (Color online) Room temperature emission spectra of all thesamples at a current injection of 20 mA. For comparison, the PL peak
positions from n-side active regions and p-side active regions are indicatedby the solid and dashed arrows, respectively.

diﬀerent from the PL spectra. For comparison between theemission peak positions by current injections and the PL
peak positions shown in Fig. 3, the PL peak positions arealso indicated in Fig. 4 by the solid (for n-side active) andthe dashed (for p-side active) arrows. The sample with the
0 nm Mg-doped layer showed only a blue emission from thep-side active region. This seems reasonable because holes
are only injected into the p-side active region due to muchheavier mass and shorter diﬀusion length of the holes. Thenthe samples with the 0.8 and 1.6 nm Mg-doped layers clearly
showed a violet and a blue emissions simultaneously fromthe n- and p-side active regions. On the contrary, the samples
with the 4 and 14 nm Mg-doped layers showed only a violet
emission from the n-side active regions. These results
suggest that appropriate Mg doping at the intermediate layer,
corresponding to the cases of the 0.8 and 1.6 nm Mg-doped
layers, should supply enough holes even into the n-side
violet active region. In other words, distributions of holes
and electrons are well-balanced in both the n- and p-side
active regions, so that two reasonable peak intensities wereobtained. On the other hands, more Mg doping, such as 4
and 14 nm Mg-doped layer cases, should not only enhance ahole injection into the n-side violet active region but also
suppress an electron injection into the p-side blue active
region, resulting in a single violet emission. These results
indicate that the balance of the emission intensities between
the two active regions can be designed by the Mg doping
level at the intermediate layer.
Figure 5 shows I–L characteristics of all the samples.
We found that the light output intensities of the samples
with the 0.8, 1.6, and 4 nm Mg-doped layers were less thanhalf of those of the samples with 0 and 14 nm Mg-doped
layers. This result most likely came from an inﬂuence of theunintentional Mg incorporation into the p-side active regions
shown in Fig. 2. Since Mg acts as a nonradiative center, thenthe radiative eﬃciency of the p-side active region should bedecreased. At this moment our explanation for the above
result is as follows. Obviously the sample with 0 nm
Mg-doped layer was not inﬂuenced by the unintentional
Mg incorporation, resulting in the strong emission from thep-side active region. In the case of the 14 nm Mg-doped
layer, not only the intermediate layer but also the p-side
active region are probably p-type due to the unintentional
Mg incorporation, so most electrons and holes are eﬃciently
injected into the n-side active region. On the contrary, thesamples with 0.8, 1.6, and 4 nm Mg-doped layers could
contain moderate Mg concentrations in the intermediate
layer, so electrons and holes are injected to both the p- andn-side active regions. Then some carriers in the p-side active
regions recombine through the Mg-related nonradiative

08JG02-3# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JG02

K. Matsui et al.

ing two active regions with the intermediate layer, while thatof our conventional LED were 3.33 V. The high forward
voltages for the samples with the two active regions arepresumably due to the thicker total thicknesses (135 nm
instead of 75 nm in the conventional LED) of the active
regions and also undoped barriers of the MQWs.

p-sideNormalized intensity[a.u.]n-side
100 mA

4. Summary

50 mA

20 mA

350

400

450

500

550

Wavelength[nm]Fig. 6. (Color online) Current injection dependence of emission spectraof the sample with the 1.6 nm Mg-doped layer.

Voltage[V]5
4

Acknowledgments

3
2
1
0
0

We have grown the two diﬀerent active regions separated
with the Mg-doped intermediate layer placed in the single
pn-junction. Room temperature PL spectra of all the samples
basically showed similar spectrum patterns, meaning that theactive region qualities were similar. At the same time, very
diﬀerent emission spectra by current injection of 20 mA
were observed among the samples. The samples with the
0.8 and 1.6 nm Mg-doped layers in the intermediate layers
clearly showed the two diﬀerent emissions simultaneously
from the n- and p-side active regions. At the same time, thesamples showing two emissions resulted in lower light
output intensities by current injection, possibly caused by theunintentional Mg incorporation in one of the active regions.
The intensity ratio between the two emissions was changed
by changing the current injection.

This work was partly supported by the Hibi Science
Foundation and Ministry of Education, Culture, Sports,
Science and Technology Supported Program for theStrategic Research Foundation at Private Universities,
2012–2016. The SIMS measurements were carried out bycourtesy of Toray Research Center, Inc.

14 nm
4 nm
1.6 nm
0.8 nm0 nm
Conventional LED

5 10 15 20 25 30 35 40 45 50
Current[mA]

Fig. 7. (Color online) Room temperature I–V characteristics for all thesamples and a conventional LED.

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all the injected carriers in the p-type active region recombine
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08JG02-4

and M. Hilpert: Appl. Phys. Lett. 74 (1999) 1367.

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Concentrating Properties of Nitride-Based Solar Cells Using Different Electrodes

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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JH02
http://dx.doi.org/10.7567/JJAP.52.08JH02

Concentrating Properties of Nitride-Based Solar Cells Using Different Electrodes
Mikiko Mori1 , Shinichiro Kondo1 , Shota Yamamoto1 , Tatsuro Nakao1 , Motoaki Iwaya1 ,Tetsuya Takeuchi1 , Satoshi Kamiyama1 , Isamu Akasaki1;2 , and Hiroshi Amano2;3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
2

Received October 11, 2012; accepted December 14, 2012; published online May 20, 2013
We investigated the concentration properties of GaInN-based solar cells using different window electrodes. A significant difference was observed
between the concentrating properties of the window electrode structures. It was clearly found that indium tin oxide (ITO) is suitable as anelectrode. The short-circuit current density, open-circuit voltage, fill factor, and energy conversion efficiency of nitride-based solar cells fabricatedusing an ITO electrode were 7:1  102 mA/cm2 , 2.2 V, 79%, and 4.0%, respectively, under an air mass filter of 1.5G at 300 suns and at room
temperature. # 2013 The Japan Society of Applied Physics

1. Introduction

Nitride material systems including AlN, GaN, InN, and their
alloys are widely used as light-emitting diodes and laser
diodes.1–4) Since the band gap of GaInN ternary alloys
covers a broad range from 0.65 to 3.43 eV,5) these alloys are
suitable for solar cell applications.6–8) Up to now, a few
studies on III–nitride solar cells have been reported.9–11) We
have previously reported the fabrication of GaInN-basedsolar cells.12–16) By improving the crystal quality of GaInN
with a low InN molar fraction by applying GaInN
superlattice (SL) structures,13) the conversion eﬃciency of
nitride-based solar cells has been improved. We also
reported that it is essential to realize a low pit density,
which should be lower than 107 cm2 , to realize highperformance GaInN-based solar cells with a high opencircuit voltage (VOC ) and ﬁll factor (FF) at 1 sun.14)
However, when GaInN-based solar cells with a high InN
molar fraction were fabricated, favorable performance was
not observed because of the increase in the leakage current.
Furthermore, there are several disadvantages of nitridebased solar cells. One of the disadvantages of compound
semiconductor solar cells is their high cost compared with
Si and organic solar cells. This problem can be solved byusing a solar concentrator. Furthermore, a condenser can be
designed to lower the production cost. The performances ofthe solar cells using other material systems such as AlGaInPand GaInAs have been improved by the concentration ofsunlight.17–20) Although we have reported the properties ofnitride-based solar cells under a concentrated light illumination,16) they have not been analyzed in detail. In particular,
the window electrode pattern under concentration light has
not been discussed in detail. In nitride semiconductor solar
cells, the resistivity of p-GaN is high, making it diﬃcult toform a satisfactory ohmic contact with the material. In
addition, there are few metals with a high work function
comparable to the work function of p-GaN that can also
form a low-resistance interface with GaN. The structureof most nitride-based solar cells reported so far has beenp-side-up, because it is diﬃcult to realize a p-type ohmic
contact when p-GaN is etched by dry etching. In general, theresistivity of p-GaN is two orders of magnitude higher thanthat of n-GaN. Therefore, the p-electrode of nitride-basedsolar cells may be an essential feature of the high current

Fig. 1. (Color online) Schematic view of device structure.

spreading. Generally, a grid electrode is used in solar cells
based on other material systems. Probably, there is a suitableelectrode structure for each material. It is important toinvestigate the electrode structure to improve the conversion
eﬃciency of nitride-based solar cells.
In this study, we optimized the window electrode pattern
for GaInN-based solar cells. We also investigated theconcentrating properties of such nitride-based solar cells atlight intensities of up to 300 suns.
2. Experimental Methods

Figure 1 shows a schematic view of the device structure.
The device structure is a GaInN-based solar cell with a
GaInN/GaInN superlattice active layer on a sapphire
substrate covered with a low-temperature buﬀer layer.21)
We grew 50 pairs of unintentionally doped Ga0:83 In0:17 N(3 nm)/Ga0:93 In0:07 N (0.6 nm) superlattice layers as active
layers on freestanding GaN substrates. Another 10 pairsof Si-doped Ga0:90 In0:10 N (3 nm)/GaN (3 nm) superlatticelayers were inserted beneath the active layers. The Siconcentration in the 10 pairs of superlattice layers was
3  1018 cm3 . The role of the additional Si-doped superlattice layers was to suppress the generation and growth ofV-shaped pits.22) The density of V-shaped pits in this sample
is lower than 1  107 cm2 , which is suﬃcient to achieve

08JH02-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JH02(a)

M. Mori et al.

(b)

(d)
(c)
Fig. 2. (Color online) (a)–(c) Schematic views of electrode structure: (a) Ni/Au, (b) Ni/Au grid, and (c) ITO. (Orange line is auxiliary line of contact
area.) (d) Size of the grid electrode in detail.

100
80

Transmittance [%]high-performance nitride-based solar cells with high VOC
and FF at 1 sun.
Figures 2(a)–2(c) show microscopy images observed
from the top of each device. Figure 2(d) also shows thedimensions of the Ni/Au grid electrode in detail. Threesamples with diﬀerent p-electrodes were investigated. In
the type I sample [Fig. 2(a)], Ni/Au whole-area electrodes(5 nm/5 nm) were deposited on a p-GaN layer by electron
beam evaporation, and rapid thermal annealing was
performed at 525  C in oxygen for 5 min. This electrodestructure is often used in typical nitride-based LEDs. In thetype II sample [Figs. 2(b) and 2(d)], Ni/Au (5 nm/5 nm)
grid electrodes were deposited on p-GaN layers by electron
beam evaporation. This structure is often used in solar cells
based on other material systems. Details of the electrodestructure such as the sizes of ﬁnger and bus-bar electrodes
are shown in Fig. 2(d). In the type III sample [Fig. 2(c)],indium tin oxide (ITO) whole-area electrodes (100 nm) were
deposited on p-GaN layers by sputtering, and rapid thermal
annealing was performed at 550  C in nitrogen for 10 min.
An ITO electrode has commonly been used in recent blue
LEDs. Ti/Al/Ti/Au (30 nm/100 nm/20 nm/150 nm) electrodes were deposited by electron beam evaporation as
ohmic contacts to the n-GaN layers.
Figure 3 shows the transmittances of the Ni/Au and ITOelectrodes. Table I also summarizes the transmittances ofthese electrodes and the speciﬁc contact resistance between
p-GaN and these electrodes.23) Although the transmittance of
the ITO electrode in the visible region is higher than that of
the Ni/Au electrode, the contact resistance and speciﬁc
contact resistance between p-GaN are higher than those for
the Ni/Au electrode. In addition, the transmittance of the
ITO electrode in the UV region (300 nm or less) is lower
than that of the Ni/Au electrode. In this study, the structureof the Ni/Au grid electrode and the thickness of the ITO
electrode were optimized at 1 sun.
The conversion eﬃciencies of these devices were
measured using an air mass 1.5 G (AM 1.5G) solar simulator(Asahi Spectra HAL-320). Both the irradiation area and thelight intensity were varied to concentrate the light. The
current density versus voltage characteristics of each device
were measured at light intensities from 1 to 300 suns at room

60
40
20
0

ITO(100nm)
Ni/Au(5nm/5nm)

300

400
500
600
Wavelength [nm]700Fig. 3. (Color online) Transmittances of ITO and Ni/Au electrodes.

Table I. Characteristics of Ni/Au and ITO electrodes.

Material

Thickness Sheet resistance
(nm)
(/)

Ni/Au
10
(5 nm/5 nm)
ITO
1:0  102
(100 nm)Speciﬁc contact Transmittance(400 nm)resistance(%)(cm2 )14

5:3  10458

2:9  102

2:6  103

92

temperature (RT). We investigated the concentrating properties of such nitride-based solar cells at light intensities of up
to 300 suns.
3. Results and Discussion

Figure 4 shows the current density versus voltage characteristics of the nitride-based solar cells under the solar
simulator (AM 1.5G, 1 to 300 suns) at RT. The solar cells
exhibited good characteristics from 1 to 300 suns. We foundthat JSC and VOC for the nitride-based solar cells increased
with increasing concentration ratio.
The short-circuit current density (JSC ), VOC , series
resistance (RS ), shunt resistance (RSH ), FF, and energy
conversion eﬃciency () at 1 and 300 suns for each sample

08JH02-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JH02

2JSC [mA/cm ]M. Mori et al.

800
700
600
500
400
300
200
100
0
3.0

Ni/Au
Ni/Au grid
ITO

VOC [V]

2.5
2.0
1.5
1.0
0

50 100 150 200 250 300
Concentration Ratio

Fig. 5. (Color online) JSC and VOC as a function of concentration ratio.

Fig. 4. (Color online) I–V curves of solar cells using each electrode
under solar simulator (AM 1.5G, 1 to 300 suns) at RT.

Table II. Device performance characteristics under solar simulator (AM
1.5G, 1 sun) at RT.

p-electrode

VOC
(V)

FF
(%)

JSC
(mA/cm2 )


(%)

RS
(cm2 )RSH(cm2 )

Ni/Au

1.8

82

1.0

1.5

1:7  102

3:4  104

1.9

2

4:9  10

3:9  104

2.2

2

3:4  104

Ni/Au grid
ITO

2.0
1.8

82
82

1.2
1.5

2:1  10

Table III. Device performance characteristics under solar simulator (AM
1.5G, 300 suns) at RT.

p-electrode

VOC
(V)

FF
(%)

JSC
(mA/cm2 )


(%)

RS
(cm2 )RSH(cm2 )Ni/Au
Ni/Au grid

2.1
2.2

81
64

4:5  102
3:7  102

2.5
1.8

1.0
3.2

34
85

ITO

2.2

79

7:1  102

4.0

0.8

50

using diﬀerent electrodes are summarized in Tables II and
III, respectively.
Figure 5 shows JSC and VOC as a function of theconcentration ratio. We compared these results with thesolar cell characteristics of other material systems. In theory,
JSC for solar cells should increase linearly with increasing
concentration ratio. From Fig. 5, JSC for all samples

increased linearly with increasing concentration ratio.
Among the three types of electrode structure, JSC is highest
for the ITO electrodes. This is because the transmittance of
ITO electrodes is higher than that of Ni/Au electrodes.
Next, we consider VOC , where VOC for a condenser can bedeﬁned as follows:24)


nkT
JSC
ln
VOC ¼
þ1 :
ð1Þ
q
J0
Here, J0 , n, k, T , and q are the reverse saturation current,
nonideality factor, Boltzmann constant, temperature, andelementary charge, respectively. According to Eq. (1), VOC
should increase logarithmically with the increased condensations of light. The results for all samples almost match thetheory from 1 to 300 suns. Moreover, among the three typesof electrode structure, VOC for the solar cell using a Ni/Au
grid electrode was slightly higher. This slight diﬀerence iscaused by the small coverage of pits by the Ni/Au gridelectrode, because the contact area of the Ni/Au gridelectrode is narrower than that of the other electrodes.
Figure 6 shows FF and  as a function of the concentration ratio. Although FF for the solar cells using ITO and Ni/
Au whole-area electrodes remained constant upon changing
the concentration ratio, that of the solar cell using a Ni/Au
grid electrode was greatly reduced with increasing concentration ratio. As a result,  for the solar cell using the Ni/Au
grid electrode showed almost no increase with increasing
concentration ratio. We next discuss the cause of thediﬀerence in the concentrating characteristics of FF in eachsolar cell.
Figure 7 shows RS and RSH for each solar cell as afunction of the concentration ratio. Although RS and RSH forall samples decrease with increasing concentration ratio, therate of decrease of RS in the solar cells using a Ni/Au grid

08JH02-3# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JH02

M. Mori et al.

0.85
0.80

FF

0.75
0.70
0.65
Ni/Au
Ni/Au grid
ITO

0.60
Conversion Efficiency [%]0.55
5Fig. 8. Equivalent circuit of solar cell. Iph and Rload are photocurrent andexternal resistance, respectively.

4
3
2
1
0

50 100 150 200 250 300
Concentration Ratio

Fig. 6. (Color online) FF and  as a function of concentration ratio.

2

RS [Ω cm ]

1000

Ni/Au
Ni/Au grid
ITO

100

4. Conclusions

We investigated the correlation between the device performance and the window electrode pattern for GaInN-basedsolar cells. To realize high-performance GaInN-based solar
cells with a high VOC and FF, it is essential to realize a low
RS and high RSH . VOC and  for the cells were found toincrease with increasing light intensity. The conversion
eﬃciency of the solar cell using an ITO p-electrode was thehighest at light intensities from 1 to 300 suns because of its
good performance in increasing the concentration of incident
light and its low RS . As a result, we obtained  of up to 4.0%
by irradiating concentrated sunlight with a light intensity of
up to 300 suns.

10
1

1000002

RSH [Ω cm ]where IC is the circuit current ﬂowing to Rload . Therefore,
the extraction of energy from the solar cell using the Ni/Au
grid electrode is decreased with increasing circuit current
owing to the high concentration ratio. Since the resistivity ofthe p-type nitride semiconductor layer is higher than that ofother material systems, we concluded that the grid electrodeis not suitable for GaInN-based solar cells with the p-side-up
structure. Therefore, we concluded that the ITO whole-area
electrode is suitable as an electrode for GaInN-based solar
cells with the p-side-up structure, although there is an issue
regarding the cost due to the scarcity of indium.26) We also
believe that in the future there will be a need to consider
GaInN-based solar cells with the n-side-up structure27,28)
fabricated using the laser liftoﬀ method29,30) and othermethods. To realize high-performance GaInN-based solar
cells, it necessary to achieve good current spreading of thep-electrode. VOC , JSC , FF, and  for the solar cell with the
ITO whole-area contacts were 2.2 V, 7:1  102 mA/cm2 ,79%, and 4.0% at 300 suns, respectively.

10000
1000
100

Acknowledgments

10
0

50 100 150 200 250 300
Concentration Ratio

Fig. 7. (Color online) RS and RSH as a function of concentration ratio.

electrode is lower than that of other samples. The highresistance of p-GaN may aﬀect the solar cell using theNi/Au grid electrode. From the equivalent circuit of thesolar cell shown in Fig. 8,25) the energy loss (Ploss ) due to theseries resistance of the solar cell can be represented asfollows:
Ploss ¼ IC2 RS ;

ð2ÞThis study was partially supported by the New Energy andIndustrial Technology Development Organization (NEDO)
Project, ‘‘Research and Development on Innovative Solar
Cells’’ and the Program for the Strategic Research Foundation at Private Universities, 2012–2016, supported by the
Ministry of Education, Culture, Sports, Science andTechnology, Japan (MEXT).

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08JH02-5# 2013 The Japan Society of Applied Physics

17)
18)
19)
20)
21)
22)

23)

24)
25)
26)
27)
28)
29)

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My IOPscience

Low-ohmic-contact-resistance V-based electrode for n-type AlGaN with high AlN molar
fraction

This content has been downloaded from IOPscience. Please scroll down to see the full text.
2016 Jpn. J. Appl. Phys. 55 05FL03(http://iopscience.iop.org/1347-4065/55/5S/05FL03)
View the table of contents for this issue, or go to the journal homepage for more
Download details:
IP Address: 219.127.71.134
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layer
Sang Min Jung, Chang Taek Lee and Moo Whan ShinREGULAR PAPER

Japanese Journal of Applied Physics 55, 05FL03 (2016)
http://doi.org/10.7567/JJAP.55.05FL03

Low-ohmic-contact-resistance V-based electrode for n-type AlGaN
with high AlN molar fraction
Kazuki Mori1*, Kunihiro Takeda1, Toshiki Kusafuka1, Motoaki Iwaya1, Tetsuya Takeuchi1,
Satoshi Kamiyama1, Isamu Akasaki1,2, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-0073, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan2*E-mail: 143434035@ccalumni.meijo-u.ac.jp
Received December 29, 2015; accepted February 23, 2016; published online April 25, 2016
We investigated a V-based electrode for the realization of low ohmic-contact resistivity in n-type AlGaN with a high AlN molar fraction characterized
by the circular transmission line model. The contact resistivity of n-type Al0.62Ga0.38N prepared using the V/Al/Ni/Au electrode reached1.13 ' 10%6 Ω cm2. Using this electrode, we also demonstrated the fabrication of UV light-emitting diodes (LEDs) with an emission wavelength ofapproximately 300 nm. An operating voltage of LED prepared using a V/Al/Ni/Au electrode was 1.6 V lower at 100 mA current injection than that
prepared using a Ti/Al/Ti/Au electrode, with a speciﬁc contact resistance of approximately 2.36 ' 10%4 Ω cm2 for n-type Al0.62Ga0.38N.
© 2016 The Japan Society of Applied Physics

1.

n-Al0.62 Ga0.38 N (2.0μm)

Introduction

Ultraviolet (UV) light has been used in various ﬁelds, suchas medicine, chemical treatment, light sensing, and semiconductor manufacturing processes. AlGaN-based deep-UV
light-emitting diodes (UV LEDs) are expected to be usedas the next-generation UV-light emitters that will replace
conventional UV lights such as mercury lamps. UV LEDs
have many advantages over mercury lamps, such as theirsmall size, long lifetime, and high eﬃciency. Currently, highexternal quantum eﬃciencies of UV LEDs ranging between UV-B (emission wavelength 280–315 nm) and UV-C
(emission wavelength: lower than 280 nm) have already beenreached.1–9) From a practical viewpoint, improvement of wall
plug eﬃciency (WPE) of these LEDs is essential. A decreasein device resistance, including the sheet and speciﬁc contact
resistances (ρc) of n- and p-type layers, is important for
WPE improvement. In particular, the ρc of the n-type AlGaN
layer with a high AlN molar fraction accounts for a large
component of device resistance in these LEDs. Thus, areduction in ρc in the n-type AlGaN layer with a high AlN
molar fraction is important.
Many groups have reported the fabrication of an electrodefor n-type GaN with good ohmic contact using Ti-, Mo-,
and Cr-based four-metal-layer structures.10–12) The ﬁrst-layer
metal in this electrode is Ti, Mo, or Cr. The metal reacts with
GaN to form a nitride metal and generates N vacancies in
GaN as donors, thereby increasing the doping level at themetal–semiconductor interface. The second layer is almost
always Al, which ensures that the aforementioned dopinglevel increase does not occur by impeding the out-diﬀusionof Ga.13,14) The third layer is a barrier layer that stabilizes theelectrode and prevents second-layer oxidation and melting.
The fourth layer is a cap layer that prevents the oxidation of
the underlying layers.
Ti-based electrodes can form good ohmic contact withn-type AlGaN with an AlN molar fraction of up to 0.5.15–19)
However, Ti-based electrodes increase speciﬁc contact
resistance, thereby increasing the AlN molar fraction. Tibased electrodes have high speciﬁc contact resistance ton-type AlGaN with AlN molar fractions above 0.6.20)

AlN (2.5μm)
Sapphire substrate
Fig. 1. Schematic structure of n-Al0.62Ga0.38N.

Recently, France et al. reported a low speciﬁc contact
resistance of n-type AlGaN using a V=Al=V=Au electrodewith an AlN molar fraction of 0.6.21) However, we havefound no report on the application of this electrode for UV
LEDs, because this electrode has low repeatability. The
V=Al bilayers are unreliable after high-temperature annealing
because of the propensity for V=Al oxidation to occur and thetendency of Al to melt and ball up. Thus, the barrier layer isimportant for improving the thermal stability of contacts.
In this study, we investigated the barrier layer of a V-based
electrode to improve its stabilization against high-temperature annealing. We evaluated the contact characteristics
of each barrier layer under various annealing conditions.
Moreover, we demonstrated the fabrication of UV LEDs with
300 nm emission wavelengths using this electrode.
2.

Experimental procedure

Figure 1 shows a schematic view of the sample structure. The
samples were epitaxially grown on a c-plane sapphire
substrate by metal–organic vapor phase epitaxy. An unintentionally doped, 2.5-µm-thick AlN layer was grown at
1,180 °C and a 2.0-µm-thick Si-doped Al0.62Ga0.38N layer
was sequentially grown. The free-electron concentration in
Si-doped Al0.62Ga0.38N was determined to be approximately
1.9 × 1018 cm−3 at room temperature (RT) by Hall-eﬀect
measurements using the van der Pauw method. Moreover,
the resistivity and mobility of this Al0.62Ga0.38N were 6.9 ×
10−2 Ω cm and 49 cm2 V−1 s−1, respectively.
After cleaning with an organic solvent, aqua regia, andhydrogen ﬂuoride, the V (20 nm)=Al (80 nm)=barrier layer(40 nm)=Au (100 nm) electrodes were deposited on this ntype Al0.62Ga0.38N. Thus far, a Ti-based electrode has been

05FL03-1© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FL03 (2016)

5μm

K. Mori et al.

10μm

15μm

25μm

20μm

30μm

250μm
Fig. 2. (Color online) Microscopy image of the electrode patterns for measurements.

Ti/Au pad

1.0

Ni/Au p-electrode(10nm/40nm)

0.5
Current [mA]p-GaN (100nm)
p-Al0.6 Ga0.4 N (150nm)
Al0.45 Ga0.55 N (3.0nm)/
Al0.55 Ga0.45 N (9.0nm) 3QWs

n-electroden-Al0.6 Ga0.4 N (2.0μm)
AlN (2.5μm)

V/Al/V/Au
V/Al/Mo/Au
V/Al/Ni/Au
V/Al/Ti/Au
Ti/Al/Ti/AuElectrode interval
20μm

-0.5

0.0

0.5

1.0

Fig. 4. (Color online) I–V characteristics of each electrode at an electrodeinterval of 20 µm.

n-electrodep-electrode
500μm
(b)

Fig. 3. (Color online) (a) Structure of UV LEDs with an emitting peakwavelength of 300 nm and (b) Plane view of the electrode pattern.

reported to be introduced into the barrier layer (Ni, Ti, andMo) between the Ti=Al and Au layers. Therefore, weinvestigated the dependence of barrier layers on V, Mo,
Ti, and Ni. After electrode deposition, the samples wereannealed for various times from 1 to 7 min under ambient N2
at 900 °C. We also prepared a sample using a Ti=Al=Ti=Au
electrode as a reference.
Contact characteristics were measured using the circulartransmission line model (CTLM) at RT.22) Figure 2 shows
the CTLM patterns obtained in this study. The electrodeintervals were 5, 10, 15, 20, 25, and 30 µm. The propagation
length (Lt) and sheet resistance (RSH) were calculated fromthis measurement. Therefore, ρc was derived fromc ¼ RSH  L2t :-0.5

Voltage [V]

(a)

145μm

0.0

-1.0
-1.0

C-plane sapphire substrate

245μm

RT
DC

ð1Þ

The structure of UV LEDs with an emitting peakwavelength of 300 nm is shown in Figs. 3(a) and 3(b). UV
LEDs were composed of a 2.5-µm-thick AlN template, a
2.0-µm-thick n-type Al0.6Ga0.4N cladding layer, three pairsof unintentionally doped Al0.45Ga0.55N=Al0.55Ga0.45N multi-

quantum wells (3QW), a 150-nm-thick p-type Al0.6Ga0.4N
cladding layer, and a 100-nm-thick p-type GaN contact layer.
These samples were annealed at 800 °C for 5 min in ambient
N2 for p-type activation. These samples were etched up ton-type AlGaN by inductively coupled plasma reactive-ion
etching. Then, the n-type electrodes were deposited on the
n-type AlGaN layer by electron-beam evaporation. The
size of the n-type electrode was 500 × 245 µm2. After thedeposition of these electrodes, they were annealed at 900 °C
for 5 min in ambient N2 to induce ohmic contact to the n-type
AlGaN layer. Ni=Au (10=40 nm) was deposited on the p-type
GaN layer and annealed at 550 °C for 3 min in ambient O2 toinduce ohmic contact to the p-type GaN layer. The size ofthe p-type electrode was 500 × 140 µm2. In addition, Ti=Au(10=200 nm) pad electrodes were formed on all electrodes
for implementation. Injection voltage–current (V–I) characteristics were measured using a semiconductor parameter
analyzer (Agilent 4156C) at RT and DC.
3.

Results and discussion

Figure 4 shows the current–voltage (I–V) characteristics at
20 µm intervals between the Ti=Al=Ti=Au electrodes andV-based electrodes with diﬀerent barrier layers. For thisﬁgure, we describe only the result of optimizing theannealing conditions for each electrode. The optimized
annealing times of V=Al=Ti=Au, V=Al=V=Au, V=Al=Ni=
Au, and V=Al=Mo=Au electrodes were 3, 4, 5, and 5 min,respectively.
As shown in the ﬁgure, the V=Al=V=Au and V=Al=Mo=Au
electrodes exhibited nonlinear I–V characteristics. Introducing
V or Mo as a barrier layer formed a Schottky contact. In
contrast, the V=Al=Ni=Au and V=Al=Ti=Au electrodes wereshown to have good ohmic-contact characteristics. Moreover,the slopes of the I–V characteristics of the V=Al=Ni=Au and
V=Al=Ti=Au electrodes were larger than that of the Ti=Al=Ti=
Au electrode. Introducing Ni or Ti as a barrier layer formed

05FL03-2© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FL03 (2016)

Resistance (Ω)

80Resistance (Ω)

100

V/Al/Ni/Au

60
40
20
0

0

80
60
40
20
0

5 10 15 20 25 30 35
Electrode interval (μm)

0

V/Al/Mo/Au

3000
2000
1000
0

5 10 15 20 25 30 35
Electrode interval (μm)(a)

0

5 10 15 20 25 30 35
Electrode interval (μm)
(c)

(b)

300

4000

Ti/Al/Ti/Au

V/Al/V/Au

3000

Resistance (Ω)

Resistance (Ω)

4000

V/Al/Ti/Au
Resistance (Ω)

100

K. Mori et al.

2000
1000
0

0

5

200
100
0

10 15 20 25 30 35

05

10 15 20 25 30 35

Electrode interval (μm)

Electrode interval (μm)

(d)

(e)

Fig. 5. (Color online) Resistance at each electrode interval and optimized annealing time: (a) V=Al=Ni=Au, (b) V=Al=Ti=Au, (c) V=Al=Mo=Au,(d) V=Al=V=Au, and (e) Ti=Al=Ti=Au.

1

10

2

Specific contact resistance [Ωcm ]

0

10

-1

10

V/Al/Ni/Au

-2

10

30μm

30μm

V/Al/Ti/Au
-3

10

(b)

(a)

V/Al/Mo/Au
V/Al/V/Au-4

10

Ti/Al/Ti/Au

-5

10

-6

10

0

1

2
3
4
5
6
Annealing time [min]

7

30μm

(c)

Fig. 6. (Color online) ρc of each barrier layer depending on the annealing
conditions.

30μm

(d)

Fig. 7. Diﬀerential interference microscopy images of (a) V=Al=Ni=Au,(b) V=Al=Ti=Au, (c) V=Al=Mo=Au, and (d) V=Al=V=Au after annealing.

better ohmic contact to the V=Al=Ni=Au and V=Al=Ti=Au
electrodes than to the Ti=Al=Ti=Au electrode.
For calculating ρc, resistances between electrodes were
measured at each electrode interval. Figures 5(a)–5(e) show
the resistance plots at the electrode interval from 5 to 30 µm.
Since the resistance plots almost agree with the ﬁtting line,
the speciﬁc contact resistances obtained are considered
reliable in this investigation.
Figure 6 shows the annealing-time dependence of ρc ateach electrode at 900 °C in ambient N2. The ρc values of the
V=Al=V=Au and V=Al=Mo=Au electrodes were higher than
2.84 × 10−2 and 2.73 × 10−2 Ω cm2, respectively, even when
the annealing time was changed from 1 to 7 min. This is dueto the Schottky contact. In contrast to each electrode, the ρcvalues of V=Al=Ti=Au and V=Al=Ni=Au electrodes wereas low as 2.90 × 10−6 and 1.13 × 10−6 Ω cm2, respectively.
We found that these two electrodes exhibit low contact
resistances for deep UV LEDs, because the ρc value of the
Ti=Al=Ti=Au electrode was 2.36 × 10−4 Ω cm2 for an annealing time of 5 min. Therefore, we found that Ni and Ti arepromising barrier metals for V-based electrodes.
Next, we investigated the mechanism under lying the eﬀectof each barrier metal by the diﬀerential interference microscopy method and Auger electron spectroscopy (AES).
Figures 7(a)–7(d) show the diﬀerential interference microscopy images of each electrode after annealing. We found

05FL03-3© 2016 The Japan Society of Applied Physics

K. Mori et al.

60

12

50

10
N
O
Al
V
Ga
Au
Ni

30
20
10
0

0

100
200
300
film thickness [nm]Ti/Al/Ti/Au
V/Al/Ni/Au

8
6

Intensity [arb.unit.]

40

Voltage [V]atomic concentration [at%]Jpn. J. Appl. Phys. 55, 05FL03 (2016)

4
2

250

0

400

0

20

atomic concentration [at%]

(a)

40

275
300
325
Wavelength [nm]60

Current [mA]

80

350

100

Fig. 9. (Color online) V–I characteristics of UV LEDs emitting at 300 nm
prepared using Ti=Al=Ti=Au and V=Al=Ni=Au n-type contacts.

60
50
40
N
O
Al
V
Ga
Au
Ni

30
20
10
0

0

100
200
300
film thickness [nm]

400

(b)
Fig. 8. (Color online) Atomic concentrations analyzed by AES:(a) V=Al=V=Au and (b) V=Al=Mo=Au.

a signiﬁcant diﬀerence in surface morphology of eachelectrode. The surface roughness of the electrode fabricatedusing V and Mo barrier metals was much larger than thatfabricated using Ni and Ti barrier metals. In particular, wefound protruding portions on the surfaces of both samples,
which were conﬁrmed by the laser microscopy method. We
also measured the atomic concentration by AES of two
samples fabricated using Ni and V barrier metals. These two
samples show signiﬁcant diﬀerences in surface roughness
and contact resistance. Figures 8(a) and 8(b) show the depth
proﬁles of the V=Al=Ni=Au and V=Al=V=Au electrodes.
From these ﬁgures, although the outermost surface was Au
before thermal annealing, the V=Al=Ni=Au electrode untila depth of 50 nm from the surface was composed almost
entirely of Al and O, i.e., the V=Al=Ni=Au electrode wasformed with an AlOx oxide layer as a result of thermal
annealing. The Ni–Al alloy formed by thermal annealing
conﬁned the stable oxide ﬁlm and the control of oxidation to
the underlying metals.23) Thus, the oxygen concentration at
the contact interface was 8.14%. We initially thought that
this result reﬂected incipient layer composition. However, the
V=Al=V=Au electrode is composed almost entirely of Al andO, with only small amounts of Au and Ga at the surface.
Thus, the AlOx oxide layer is degraded by high-temperatureannealing. Also, the VOx oxide layer is formed by thereaction between V and O at the barrier layer. VOx induces
the high-temperature corrosion phenomeon.24) Thus, themetal composition exhibits notable interdiﬀusion becauseof it, the underlying layers of the electrode were rapidly
oxidized. The oxygen concentration of the electrode was
17.0% at the contact interface. For high-temperature anneal-ing, the barrier layer in the V-based electrode is importantfor meeting the requirement for the control of the degree ofoxidation.
We also fabricated UV LEDs with an emission wavelengthof 300 nm with V=Al=Ni=Au as the n-type electrode, which
exhibited a ρc value of approximately 1.13 × 10−6 Ω cm2 forthe n-type Al0.62Ga0.38N. We prepared the same LEDs using
a Ti=Al=Ti=Au electrode, which exhibited a ρc value ofapproximately 2.36 × 10−4 Ω cm2 for the n-type Al0.62Ga0.38N. Figure 9 shows the V–I curves of each LED. In
this ﬁgure, we inserted the typical spectrum of this UV LED.
According to this ﬁgure, the operating voltage of the LED
prepared using a V=Al=Ni=Au electrode was reduced by
1.6 V at 100 mA current injection, compared with that
prepared using a Ti=Al=Ti=Au electrode.
4.

Conclusions

In this study, we investigated the V-based electrode for low
ρc in n-Al0.62Ga0.38N. We found it advantageous to fabricate
the V-based electrode with low ρc using Ni or Ti as a barrier
metal. The minimum ρc value of n-type Al0.62Ga0.38N
prepared using the V=Al=Ni=Au electrode reached 1.13 ×
10−6 Ω cm2. We also demonstrated the fabrication of a UV
LED with an emission wavelength of approximately 300 nm
using this electrode.
Acknowledgments

This study was partially supported by the Program for
the Strategic Research Foundation at Private Universities,
2012–2016; by a MEXT Grant-in-Aid for Specially Promoted Research (No. 25000011), and by a MEXT Grant-inAid for Scientiﬁc Research A (No. 15H02019).

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RAPID COMMUNICATION

Japanese Journal of Applied Physics 53, 030305 (2014)
http://dx.doi.org/10.7567/JJAP.53.030305

Control of growth mode in Mg-doped GaN/AlN heterostructure
Tomohiro Morishita1,2*, Kosuke Sato1, Motoaki Iwaya2, Tetsuya Takeuchi2, Satoshi Kamiyama2, and Isamu Akasaki2,3
1

Asahi Kasei Corporation, Fuji, Shizuoka 416-8501, Japan
Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
E-mail: morishita.tc@om.asahi-kasei.co.jp
2

Received November 18, 2013; accepted December 23, 2013; published online February 14, 2014
We investigated the effect of carrier gas on the growth mode of Mg-doped GaN on AlN layers. The growth mode of Mg-doped GaN strongly
depends on the type of carrier gas used. In H2 carrier gas, Mg-doped GaN on AlN layers is grown in the shape of an island. In contrast, in N2 carriergas, Mg-doped GaN on AlN layers shows a two-dimensional growth different from the initial growth mode. We observed that the type of carrier gas
used has a signiﬁcant impact on the electrical conductivity of Mg-doped GaN owing to the difference in surface ﬂatness.
© 2014 The Japan Society of Applied Physics

A

lGaN ternary nitride semiconductors show great
potential for optoelectronic applications, particularly in deep UV (DUV)/UV light-emitting diodes(LEDs). These devices have many applications such as in air/water sterilization and can be used as a substitute of mercury
lamps from the viewpoint of environmental burden. However, the external quantum efﬁciency (EQE) of DUV/UVLEDs is much lower than that of blue LEDs.1,2) There aremany studies on the increase in internal quantum efﬁciency(IQE) by reducing the dislocation density3–5) such as thatusing free standing AlN substrates, an epitaxial lateralovergrowth method, and so on.6–8) Because the relatively
high IQE of the AlGaN active layer in DUV/UV LEDs
has already been realized by these studies, light extraction
efﬁciency (LEE) is one of the most important factors for the
improvement of EQE in DUV/UV LEDs.
The LEE of ﬂip-chip visible LEDs can be improvedby using high-reﬂectivity electrodes based on Ag9) or Rh10)
on the back surface of the LEDs. Furthermore, the sapphire
substrate is removed by laser lift-off (LLO)11) and ﬂip-chip
visible LEDs are sealed with resin or silicone.12) By combining these methods, the LEE of visible LEDs reached³80%.13) Thus far, we have reported highly reﬂective p-type
electrodes formed by combining indium ITO and Al,14)
which are highly reﬂective in the UV region. We have also
reported the application of LLO to UV LEDs.15) However,
the most signiﬁcant issue is the absorption of light by ap-type GaN layer for the improvement of LEE in DUV/
UV-LEDs. To decrease the absorption of light by a Mg-doped
p-type GaN layer, the thickness of the p-type GaN layer must
be reduced or the layer must be replaced with a high-AlNmolar-fraction p-type AlGaN layer. However, this method isvery difﬁcult. The resistivity of the p-type AlGaN layer
increases with AlN molar fraction, because the activationenergy of the acceptor increases with the AlN molar fraction.16) In addition, the contact resistivity between the p-type
AlGaN layer and the electrode also increases markedly.17)
Thus, such a method is not practical because the operating
voltage of DUV/UV LEDs increases signiﬁcantly. Therefore,
Mg-doped p-type GaN is used as a p-type layer for almost
all DUV/UV LEDs.18,19) In contrast, the decrease in the
thickness of the p-type GaN layer is also difﬁcult because
GaN grown on an AlN layer shows three-dimensional
growth.20) Thus, in most DUV/UV LEDs, p-type GaN layerswith a thickness of more than 100 nm have been used,
resulting in the absorption of most light by Mg-doped p-type

Mg+ -doped GaN layer (20nm)
Mg-doped GaN layer (200nm)
AlN layer (2um)~~ Sapphire substrate ~~Fig. 1. Schematic view of the sample structure.

GaN layers. To decrease the absorption of light by a Mgdoped GaN layer, it is very important for GaN on AlN layers
to be grown in two-dimensional modes and to decrease thethickness of the Mg-doped GaN layer.
In this study, we investigated the growth mode of Mgdoped GaN on an AlN layer. As a result, we found that there
is a signiﬁcant difference in initial growth mode causedby the difference in the type of carrier gas used. Moreover,such a difference had a signiﬁcant impact on the electricalconductivity of p-GaN.
The AlN layers were grown on (0001) c-plane sapphire
substrates. After the growth of 2-µm-thick AlN layers at
1,250 °C, Mg-doped GaN layers were grown by a metalorganic vapor phase epitaxial (MOVPE) reactor. Trimethylaluminum (TMAl), trimethylgallium (TMGa), bis(ethylcyclopentadienyl)magnesium (Et-Cp2Mg), and ammonia (NH3)were used as the sources of Al, Ga, Mg, and N, respectively.
The growth temperature was controlled by direct measurements of the substrate surface using a pyrometer. Figure 1
shows a schematic view of the sample structure. The sapphire
substrates were thermally cleaned in H2 ambient at approximately 1,250 °C. Subsequently, AlN layers with a thickness of approximately 2 µm were grown at a temperatureof 1,250 °C and a pressure of 30 Torr using hydrogen (H2)carrier gas. Then, the substrate temperature was decreased
to approximately 1,030 °C and the MOVPE reactor pressure was increased to 200 Torr. Moreover, p-type Mg-doped
GaN layers with a thickness of approximately 200 nm and ahigher-Mg-concentration-doped (Mg+-) GaN contact layer
with a thickness of approximately 20 nm were grown with
H2 or N2. The V/III ratio was 1,000, which is a typical valueof the p-type Mg-doped GaN grown by MOVPE. The Mg
concentrations in Mg- and Mg+-doped GaN layers had set
the growth conditions to be 3 © 1019 and 2 © 1020 cm¹3,
respectively.

030305-1© 2014 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 53, 030305 (2014)

T. Morishita et al.

Mg-doped GaN growth/Mg+-doped GaN growth

two-dimensional growth

50 µ mthree-dimensional growth

50 µ m(a)(b)

Fig. 3. (Color online) Nomarski interference microscopic images of Mgand Mg+-doped GaN layers on AlN layers in (a) H2 and (b) N2 carrier gases.
Fig. 2. In situ reﬂectance measurement results of Mg- and Mg+-doped
GaN layers grown on AlN layers in H2 and N2 carrier gases.

To characterize the growth mode and thickness of
Mg-doped GaN layers, we used a real-time in situ opticalreﬂectometer. This reﬂectometer consisted of a semiconductor laser (­ = 650 nm) as the light source. The reﬂected light
was detected via the same probe, a semitransparent mirror,
and a photodiode. We also observed the surface morphology
by Nomarski interference and cross-sectional scanningelectron microscopies. Moreover, we characterized the
crystallinity and strain in Mg-doped GaN layers by a typical
X-ray diffraction (XRD) measurement. We measured the
Van der Pauw Hall effect with a magnetic ﬁeld of 1,000 G
to determine the hole concentration, mobility, and resistivityof each sample at room temperature (RT). A Ni/Au ohmic
electrode was deposited by electron beam evaporation on a
Mg+–GaN layer for Van der Pauw Hall effect measurement.
Note that the resistivity of the Mg+–GaN layer was ignored
because this layer was very thin and showed high resistivity.
Figure 2 shows in situ reﬂectance measurement results of
Mg-doped GaN layers grown with H2 and N2 carrier gases.
From this ﬁgure, we conﬁrmed the signiﬁcant difference
in reﬂectance between the samples. In H2 carrier gas, thereﬂectance from the Mg-doped GaN layer monotonicallydecreased. In addition, we could not conﬁrm the interference
from this ﬁgure. Therefore, it is suggested that the Mg-doped
GaN layer in H2 carrier gas shows a three-dimensional
growth mode rather than a two-dimensional growth mode. In
contrast, we conﬁrmed the interference from the reﬂection of
the Mg-doped GaN layer from the initial growth stage in N2
carrier gas. Therefore, it is suggested that the Mg-doped GaN
layer in N2 carrier gas shows a two-dimensional growth mode
different from the initial growth mode, and that the growthmode of the p-type Mg- and Mg+-doped GaN layers on AlN
layers strongly depends on the type of carrier gas used.
For further analysis, we obtained Nomarski interference
microscopic and cross-sectional SEM images of Mg-doped
GaN layers in H2 and N2 carrier gases as shown in Figs. 3
and 4, respectively. These ﬁgures revealed a signiﬁcant difference in surface structure depending on the type of carrier
gas used. In H2 carrier gas, Mg-doped GaN layers weregrown in the shape of islands. The height of each island wasapproximately 500 nm, which is approximately two times
larger than the designed value. In addition, there were no
GaN layer growth areas. On the other hand, although thesurface morphology was slightly worse, no island growth
and two-dimensional growth were observed in N2 carrier gas.

2 µm

1 µm

(a)

(b)

Fig. 4. Nomarski surface and cross-sectional SEM images of Mg- andMg+-doped GaN layers on AlN layers in (a) H2 and (b) N2 carrier gases.
The photograph was slightly tilted in order to clarify the structure in (b).

The thickness of Mg- and Mg+-doped GaN layers wasapproximately 180 nm, which is almost the same as thedesigned value. From these results, the Mg-doped GaN layer
was grown in the three-dimensional mode in H2 carrier gas
and in the two-dimensional mode in N2 carrier gas. A
different growth mode was assumed during the initial growth.
The equation for the decomposition of NH3 is
NH3 , N þ

3
H2 :
2

ð1Þ

From Eq. (1), the reaction of equilibrium moves to the lefthand side easily in H2 carrier gas; thus, it is difﬁcult todecompose NH3. In contrast, there is no difﬁculty in thedecomposition of NH3 in N2 carrier gas. This means thatthe number of N atoms contributing to the growth of the Mgdoped GaN layer at the surface varies even if the same
quantity of NH3 as the gas phase is supplied.
Hypothetical models for different growth modes are shownin Fig. 5. In H2 carrier gas, the number of N atoms at the AlN
surface is low and the migration of Ga atoms is sufﬁcient.
Therefore, the Mg-doped GaN layer was grown in the threedimensional mode after nucleation, as shown in Fig. 5(a).
To suppress the three-dimensional growth, the Mg-doped
GaN layer was grown at a lower growth temperature and ahigher V/III ratio; however, no two-dimensional growth wasobserved. On the other hand, in N2 carrier gas, the number of
N atoms at the AlN surface is sufﬁcient, there are many
opportunities for Ga atoms to interact with N atoms, and themigration of Ga atoms is suppressed. Thus, many small
crystal nuclei were generated and the initial growth of the
Mg-doped GaN layer was in the two-dimensional mode as
observed by in situ monitoring, as shown in Fig. 5(b).
For further analysis, the XRD reciprocal space mapping diffraction was obtained and isimage around the (2024)

030305-2© 2014 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 53, 030305 (2014)

T. Morishita et al.

Nucleation

Three-dimensional growth

AlN

Two-dimensional growth

AlN

AlN

(b)

(a)

Fig. 5. (Color online) Hypothetical models for growth modes of Mg-doped GaN layers in H2 and N2 carrier gases.

(a)

(b)
 (a) Mg- and Mg+-doped GaN layers in H2 carrier gas, and (b) Mg- and Mg+-doped
Fig. 6. (Color online) XRD reciprocal space mapping around (2024):
GaN layers in N2 carrier gas.

shown in Fig. 6. As shown in Fig. 6(a), the Mg-doped GaN
layer in H2 carrier gas was almost 100% relaxed in contrast
to that in N2 carrier gas, which was approximately 75%
relaxed as shown in Fig. 6(b). This is because Ga atoms
migrated easily in H2 carrier gas and the Mg-doped GaN
layer exhibited the lattice constant of GaN through lattice
relaxation.
Finally, we examined the electrical properties of the Mgdoped GaN layers grown in H2 and N2 carrier gases by using
the Van der Pauw Hall effect measurement method. For the
Mg-doped GaN layer grown in H2 carrier gas, no current
ﬂow was observed, even when the voltage was increased up
to 5 V because of island growth. In contrast, the Mg-doped
GaN layer grown in N2 carrier gas showed good electricalproperties. From the Hall measurements, the hole concen-

tration, mobility, and resistivity at RT were determined to beapproximately 5 © 1017 cm¹3, 6 cm2 V¹1 s¹1, and 1.3 ³ cm,
respectively, as shown in Table I. Moreover, the Mg concentration in this Mg-doped GaN layer characterized bysecondary ion mass spectrometry was 3 © 1019 cm¹3, and theactivation rate was approximately 1.7%, which is almost
the same as the typical value because the activation ratioof the p-type impurity (Mg) in GaN at RT is generally about
2 to 5%.21)
From the surface morphology and electrical properties,
Mg-doped GaN growth in N2 carrier gas was considered to bevery useful for forming DUV/UV-LEDs on AlN layers andthe obtained results encourage the use of thinner p-type Mgdoped GaN layers and promise an increase in the efﬁciency
of DUV/UV LEDs.

030305-3© 2014 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 53, 030305 (2014)
Table I.
gas.

T. Morishita et al.

Hall effect measurement results of Mg-doped GaN in N2 carrier

Mobility(cm2 V¹1 s¹1)

Resistivity(³ cm)

Hole concentration(cm¹3)

6

2.2

5 © 1017

In conclusion, we investigated the effect of carrier gas
on the growth of GaN layers on AlN underlying layers. The
growth mode of Mg-doped GaN strongly depends on thetype of carrier gas used. In H2 carrier gas, Mg-doped GaN
on AlN layers is grown in the shape of an island. In contrast,in N2 carrier gas, Mg-doped GaN on AlN layers shows a
two-dimensional growth mode. In addition, Mg-doped GaN
grown in N2 carrier gas shows good electrical characteristics.
Therefore, the growth of Mg-doped GaN in N2 carrier gas isvery useful for forming DUV/UV LEDs on AlN from theviewpoint of both two-dimensional growth and electricalproperties compared with the growth of Mg-doped GaN in
H2 carrier gas. These results promise an increase in the
efﬁciency of DUV/UV LEDs through a decrease in theabsorption of Mg-doped GaN.

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030305-4© 2014 The Japan Society of Applied Physics

Reduction of contact resistance in
V-based electrode for high AlN molar
fraction n-type AlGaN by using thin SiNx
intermediate layer

solidistatusphysicaPhys. Status Solidi C, 1600243 (2017) / DOI 10.1002/pssc.201600243

cwww.pss-c.com

current topics in solid state physics

Noriaki Nagata*,1, Takashi Senga1, Motoaki Iwaya1, Tetsuya Takeuchi1, Satoshi Kamiyama1,
and Isamu Akasaki1,2
12Faculty of Science and Technology, Meijo University, Nagoya 468-0073, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

Received 4 November 2016, revised 23 November 2016, accepted 23 November 2016
Published online 1 March 2017
Keywords contact resistance, D-UV LED, high AlN molar fraction, n-AlGaN* Corresponding author: e-mail 153434027@ccalumni.meijo-u.ac.jp

We found out the reduction of contact resistance in V-based
electrode for high AlN molar fraction n-type AlGaN by using
thin SiNx intermediate layer. The contact resistivity for n-type
Al0.70Ga0.30N with the V/Al/Ni/Au electrode using SiNx
intermediated layer reached 1.13  106 V cm2. Moreover,contact resistivity using SiNx intermediated layer had been
reduced more than one order of magnitude in all AlN molar

fractions from 0.62 to 0.87. Using this electrode, we also
demonstrated UV light-emitting diodes (LEDs) on n-type
Al0.70Ga0.30N underlying layer with an emission wavelength ofapproximately 283 nm. An operating LED voltage using a
V/Al/Ni/Au electrode with SiNx intermediated layer was 3.3 V
lower at 100 mA current injection than that without SiNx
intermediated layer.

ß 2017 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction AlGaN-based UV light-emitting
diodes (UV LEDs) have many advantages, such ascompactness, high efﬁciency, and wavelength selectivity,
these devices are expected to replace conventional UV-lightsources based on glass lamps. So far, a high external
quantum efﬁciency of deep UV LEDs ranging between
UV-B (emission wavelength 280–315 nm) and UV-C
(emission wavelength: lower than 280 nm) has alreadybeen reached [1–9]. However, there are several issues in UV
LEDs. Especially, an improvement in wall plug efﬁciency(WPE) in these LEDs is one of the most important issues in
AlGaN-based UV LEDs. For increase of WPE, reduction ofoperating voltage in deep UV LED is essential. The contact
resistance of n-AlGaN accounts for large component of thedevice resistance in LEDs, because low ohmic contact
resistivity electrode for n-AlGaN is difﬁcult with increase of
AlN molar fraction. Therefore, a reduction of contact
resistance in the n-type AlGaN layer with a high AlN molar
fraction is important.

Many other groups have reported the fabrication of anelectrode for n-type GaN with good Ohmic contact using Tibased, Mo-based, Cr-based, V-based four-metal-layer
structures [10–15]. The Ti-based electrodes can form good
low Ohmic contact (106 Vcm2) with n-type AlGaN withan AlN molar fraction of up to 0.4. However, Ti-based
electrodes increase speciﬁc contact resistance (rc), thereby
increasing the AlN molar fraction. In addition, Ti-based
electrodes have high rc to n-type AlGaN with AlN molar
fractions above 0.6 [16]. We reported that V/Al/Ni/Au
electrode realized low contact resistivity (106 Vcm2)
with n-type AlGaN with AlN molar fraction of 0.62 [17].
However, we conﬁrmed contact resistivity for AlGaN with
AlN molar fraction of 0.7 rapidly increased to 103 Vcm2order. Therefore, it is essential for realization of low ohmic
contact electrode with AlN molar fraction of more than 0.7.
So far, it has achieved the reduction of rc by insertion of thin
SiNx intermediate layer in AlGaN/GaN heterostructureﬁeld-effect-transistors [18]. However, there are no reports ofß 2017 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss1600243 (2 of 4)cN. Nagata et al.: Reduction of contact resistance in V-based electrode for n-type AlGaN

reduction of rc in high AlN molar fraction AlGaN single
layer by using SiNx intermediate layer.
In this study, we found out the reduction of rc in
V-based electrode for high AlN molar fraction n-type
AlGaN by using thin SiNx intermediate layer. We evaluatedthe contact characteristics of using thin SiNx layer depending on the thickness. Moreover, we demonstrated UV LEDs
with 283-nm emission wavelengths using this electrode.
2 Experimental The samples were epitaxially grown
on a c-plane sapphire substrate by metal-organic vapor
phase epitaxy. Figure 1 is a schematic view of the sample
structure. After baking the sapphire substrate at 1180 8C, an
unintentionally doped 2.5-mm-thick AlN layer was grown at
1180 8C and a 2.0-mm-thick Si-doped AlxGa1xN layer wassequentially grown. The free-electron concentrations in Sidoped AlxGa1xN were measured to be approximately
2.0  1018 cm3 at room temperature (RT) by Hall-effectmeasurements using the van der Pauw method. We prepared
the AlxGa1xN samples with different AlN molar fraction x
from 0.62 to 0.87.
After cleaning with an organic solvent, aqua regia, andhydrogen ﬂuoride, SiNx thin layer on this n-type
Al0.70Ga0.30N were deposited by electron beam evaporation
method, and the V (20 nm)/Al (80 nm)/Ni(40 nm)/Au(100 nm) electrodes were deposited by electron beam
evaporation method. We also investigated the dependenceof thickness of SiNx intermediated layers. The thickness of
the SiNx intermediated layer was adjusted by the both ofdeposition time and the deposition rate. In this experiment,
the deposition rate using the deposition conditions of
0.5Å s1. In addition, we did not ﬂow any other gases forsuppression of decompose the SiNx intermediated layer.
After electrode deposition, the samples were annealedfor 5 min under ambient N2 at 900 8C. We also prepared thesample using a Ti/Al/Ti/Au electrode as a reference.
The contact characteristics were measured using thecircular transmission line model (CTLM) at RT [19].
Figure 2 shows the CTLM patterns in this study. The
electrode intervals were 5, 10, 15, 20, 25, and 30 mm. The
propagation length (Lt) and sheet resistance (RSH) werecomputed from this measurement. Therefore, rc wasderived from

Figure 1 Schematic structure of n-AlxGa1xN.
ß 2017 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Figure 2 Schematic views of the electrode patterns formeasurements.

rc ¼ RSH  Lt 2 :ð1ÞThe structure of UV LEDs with an emitting peakwavelength of 283 nm is shown in Fig. 3. UV LEDs werecomposed of a 2.5-mm-thick AlN template, a 2.0-mm-thickn-type Al0.7Ga0.3N cladding layer, three pairs of unintentionally doped Al0.4Ga0.6N/GaN multi-quantum wells(3QW), a 150-nm-thick p-type Al0.7Ga0.4N cladding layer,
and a 100-nm-thick p-type GaN contact layer. Thesesamples were annealed at 800 8C for 5 min in ambient N2 for
p-type activation. These samples were etched up to n-type
AlGaN by inductively coupled plasma reactive ion etching.
Then, the V (20 nm)/Al (80 nm)/Ni(40 nm)/Au (100 nm)
n-type electrodes with and without SiNx intermediated layer
were deposited on the n-type AlGaN layer by electron-beam
evaporation for current injection. The size of the n-type
electrode was 500  245 mm2. After the deposition of theseelectrodes, annealing was performed at 900 8C for 5 minunder ambient N2 for inducing Ohmic contact to the n-type
AlGaN layer. Ni/Au (10/40 nm) was deposited on the p-type
GaN layer and annealed at 550 8C for 3 min under ambient
O2 for inducing Ohmic contact to the p-type GaN layer. The
size of the p-type electrode was 500  140 mm2. In addition,
Ti/Au (10/200 nm) pad electrodes were formed on allelectrodes for implementation. Injection voltage–current

Figure 3 Structure of UV LEDs with an emitting peakwavelength of 283 nm.
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Contributed
ArticlePhys. Status Solidi C (2017)

Figure 4 I–V characteristics with different thickness of SiNx
intermediated layer at an electrode interval of 20 mm.

(V–I) characteristics were measured using a semiconductor
parameter analyzer (Agilent 4156C) at RT and direct
current (DC).
3 Results and discussion Figure 4 shows thecurrent–voltage (I–V) characteristics at 20-mm intervals
between the V (20 nm)/Al (80 nm)/Ni (40 nm)/Au (100 nm)
electrodes on n-type Al0.7Ga0.3N inserted with different
thickness of SiNx intermediated layer. As shown in theﬁgure, all samples were shown to have good Ohmic-contact
characteristics. However, the slopes of the I–V characteristics in V/Al/Ni/Au electrode inserted with 3-nm-thick
SiNx intermediated layer were larger than that in the
0-nm-thick SiNx (no insertion SiNx layer). In contrast, theslope of I–V characteristics rapidly decreased, when
the SiNx thickness was increased to 5 nm. Therefore, it is
conﬁrmed that the improvement of I–V characteristics when
SiNx intermediated layers were the thin ﬁlms. Although the
optimum conditions of SiNx thickness is limited, the effectof the insertion 3-nm-thick SiNx was high reproducible
result. Therefore, we evaluated the detailed contact
characteristics in SiNx thickness of 0 (no insertion sample)
and 3 nm.
For calculating the rc, resistances between eachelectrode were measured at each electrode interval. Figure 5(a) and (b) shows the resistance plots of the electrodeinterval from 5 to 30 mm. Since the resistance plots almost

Figure 5 Resistance at each electrode interval: (a) SiNx/V/Al/Ni/
Au, (b) V/Al/Ni/Au.
www.pss-c.com(3 of 4) 1600243

Figure 6 The rc of each AlN molar fraction AlGaN determinedby the TLM method.

agree with the ﬁtting line, the rc obtained are considered
reliable in this investigation. The rc values for the V/Al/Ni/
Au electrodes with and without SiNx intermediatedlayer were 1.76  106 and 6.75  103 Vcm2, respectively.
Therefore, we succeeded the rc was reduced approximately
three order magnitudes by the insertion of the SiNx
intermediated layer between n-type Al0.7Ga0.3N and
V/Al/Ni/Au electrode. We also investigated the AlN molar
fraction dependence of the rc. Figure 6 shows the rc of each
AlN molar fraction AlGaN determined by the TLM method.
The rc had been reduced more than one order of magnitude
by insertion of SiNx intermediated layer in all AlN molar
fractions.
Next, we investigated the mechanism of the rc reduction
by insertion of SiNx intermediated layer. Figure 7(a) and (b)shows the cross-sectional transmission electron microscopic
image and Si concentration mapping image characterized byenergy dispersive X-ray spectrometry (EDX) of vicinity ofthe electrode. In order from the bottom, Al0.70Ga0.30N andthe electrode can be conﬁrmed. However, the layer of SiNx
was not conﬁrmed. Although clear diffusion into theAl0.70Ga0.30N layer was unable to verify, diffusion of Si to
the electrode was able to conﬁrm. Therefore, it will beconsidered that Si concentration of the electrode interface isFigure 7 Reaction of electrode n-Al0.70Ga0.30N interface ofvicinity of the electrode. (a) The cross-sectional transmission
electron microscopic image, (b) Si concentration mapping image
characterized by energy dispersive X-ray spectrometry (EDX).
ß 2017 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss1600243 (4 of 4)cN. Nagata et al.: Reduction of contact resistance in V-based electrode for n-type AlGaN

References

Figure 8 V–I characteristics of UV LEDs emitting at 283 nm.

increasing. So, it is estimated that the contact resistance isreduced by increasing the Si concentration in the vicinity ofthe electrode.
We also fabricated UV LEDs with emission wavelengthof 283 nm with V/Al/Ni/Au with and without 3-nm-thick
SiNx intermediated layer, respectively, as the n-type
electrode. Figure 8 shows the V–I curves of each LED.
In this ﬁgure, we inserted the typical spectra of this UV
LED. According to this ﬁgure, the operating voltage of the
LED using a V/Al/Ni/Au electrode with insertion of
3-nm-thick SiNx intermediated layer was reduced by
3.3 V at 100 mA current injection compared to that without
SiNx intermediated layer.
4 Conclusions We present novel method for realization of low ohmic contact electrode by using thin SiNx
intermediate layer. As a result, contact resistivity are
reduced from 6.75  103 to 1.76  106 Vcm2 by using
V-based electrode with SiNx intermediated layer for n-type
Al0.70Ga0.30N layer. We also conﬁrmed that reduction of
the operating voltage in 283-nm-emisssion deep UV-LED
at 100 mA was 3.3 V by insertion of SiNx intermediate
layer.
Acknowledgements This study was partially supported by
the MEXT Program for the Strategic Research Foundation atPrivate Universities (2012–2016), JSPS Grant-in-Aid for ScientiﬁcResearch A (no. 15H02019), JSPS Grant-in-Aid for ScientiﬁcResearch B (no. 26286045), JSPS KAKENHI (no. 16H06416), andJST CREST (no. 16815710).

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Nagasawa, A. Hirano, M. Ippommatsu, M. Iwaya, T.
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Amano, and I. Akasaki, Appl. Phys. Express 4, 092102(2011).
[8] M. Shatalov, W. Sun, R. Jain, A. Lunev, X. Hu, A.
Dobrinsky, Y. Bilenko, J. Yang, G. A. Garrett, L. E. Rodak,
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Combination of Indium–Tin Oxide and SiO2/AlN Dielectric Multilayer Reflective Electrodes for
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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JG07
http://dx.doi.org/10.7567/JJAP.52.08JG07

Combination of Indium–Tin Oxide and SiO2 /AlN Dielectric Multilayer Reflective Electrodes
for Ultraviolet-Light-Emitting Diodes
Tsubasa Nakashima1 , Kenichiro Takeda1 , Hiroshi Shinzato1 , Motoaki Iwaya1 , Satoshi Kamiyama1 ,Tetsuya Takeuchi1 , Isamu Akasaki1;2 , and Hiroshi Amano2;3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8062, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8062, Japan
E-mail: 123434034@ccalumni.meijo-u.ac.jp
2

Received October 13, 2012; revised January 25, 2013; accepted March 12, 2013; published online May 31, 2013
We have investigated novel reflective electrodes by combining an indium–tin oxide (ITO) layer and a SiO2 /AlN dielectric multilayer (DM) for UVlight-emitting diodes (LEDs). The reflectance of 10 pairs of SiO2 /AlN DM reached 98.5% at 350 nm. As a result, the ITO/DM electrodes
simultaneously satisfied the requirements of high reflectivity in the UV region, good contact characteristics, and current spreading. The light output
power of the UV LED with an ITO/DM electrode is 1.2 and 1.6 times higher than those of LEDs with ITO/Al and Ni/Au electrodes, respectively.
# 2013 The Japan Society of Applied Physics

1. Introduction

UV-light sources have been used in various ﬁelds, such as
medicine, biochemistry and photolithography, in semiconductor device fabrication process.1) Since AlGaN-based UVlight-emitting diodes (UV LEDs) have many advantages,
such as compactness, high eﬃciency, and wavelength
selectivity, these devices are expected to replace conventional UV-light sources based on glass lamps. Thus far, there
have been many reports on the improvement in the device
performance of UV LEDs.2–10) However, the external
quantum eﬃciency (EQE) of UV LEDs is still lower thanthose of visible LEDs.
To improve the EQE of UV LEDs, light extraction
eﬃciency (LEE) is one of the most important factors when
the internal quantum eﬃciency (IQE) of the active layer in
the UV LED is high.11) In general, the LEE of ﬂip-chip
visible LEDs can be improved using high-reﬂectivity
electrodes based on Ag or Rh on the back surface of theLEDs.12–14) However, the reﬂectivities of Ag and Rh in the
UV region are very low.15) Recently, we have reported
highly reﬂective electrodes that are combinations of indium–tin oxide (ITO) and Al.15) Although Al is highly reﬂective in
the UV region, it is diﬃcult to form an ohmic contact with
p-GaN. Therefore, we proposed a high-reﬂectance ohmic
electrode of ITO, to form good ohmic contact with p-GaN
with high transmittance, combined with highly reﬂective Al.
By using this bilayer electrode, the light output of UV-LEDs
was improved 1.2 times compared with that when using theconventional Ni/Au electrode. Because the reﬂectivity of Al
in the UV region is about 90%, further improvements indevice performance can be expected if an electrode with amuch higher reﬂectance is realized.
In this study, we propose a novel highly reﬂective
p-electrode of a SiO2 /AlN dielectric multilayer (DM)combined with ITO for UV LEDs. We fabricated andcharacterized 350 nm UV LEDs with this proposed electrode.
2. Experiments

The structure of UV LEDs with an emission peakwavelength of 350 nm is shown in Fig. 1. The samples were
epitaxially grown on sapphire substrates by metalorganic

p-electrodep-GaN(50nm)
p-Al0.20Ga0.80N(50nm)
p-Al0.40Ga0.60N(20nm)
Last barrier : u-Al0.10Ga0.90N(4nm)
Si-Al0.10Ga0.90N(8nm)/GaN(3nm):3QW

n-electroden-Al0.20Ga0.80N(800nm)
n-Al0.20Ga0.80N(3500nm)
Al0.25Ga0.75N
Al0.50Ga0.50N
AlN(3000nm)
Sapphire

Fig. 1. (Color online) Schematic diagram of the UV-LED.

vapor phase epitaxy. The UV LED structure was fabricatedon low-dislocation-density Al0:20 Ga0:75 N grown on agrooved Al0:25 Ga0:75 N/AlN template.16) The UV LED
structure was composed of a 3.5-m-thick n-Al0:2 Ga0:8 N
cladding layer, three pairs of unintentionally doped GaN(3 nm)/Si-doped Al0:1 Ga0:9 N (8 nm) multi-quantum wells(MQWs) with a 4-nm-thick unintentionally doped Al0:1 Ga0:9 N last barrier, a 20-nm-thick p-Al0:4 Ga0:6 N electronblocking layer, a 50-nm-thick p-Al0:2 Ga0:8 N cladding layer,
and a 50-nm-thick p-GaN contact layer. The samples wereannealed at 800  C for 5 min in N2 ambient for p-type
activation. A Ti/Al/Ti/Au ohmic contact with the n-AlGaN
layer exposed by reactive ion etching was deposited byelectron beam evaporation. After the evaporation, annealing
was performed at 850  C for 30 s in N2 ambient.
Then, three samples with diﬀerent p-electrodes wereprepared. In type I [Fig. 2(a)], Ni (10 nm)/Au (40 nm)
whole-area electrodes were deposited on the p-GaN layer
by electron beam evaporation. This electrode structure wasused as the control sample. In type II [Fig. 2(b)], ITO(10 nm)/Al (150 nm) whole-area electrodes were depositedon p-GaN by sputtering. Details of the fabrication conditions

08JG07-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JG07

Ti/Al n-electrode

Ti/Al n-electrode

Ti/Au

Ti/Au
Al(150nm)

Ni/Au(10/40nm)

100
80

Experimental
Simulation

Reflectance[%]500 µm140 µmITO/Al p-electrode245 µmNi/Au p-electrodeT. Nakashima et al.

ITO(10nm)

p-GaN

p-GaN
(b)

(a)

65 µm

65 µm

75 µm

145 µmTi/Al n-electrode

Ti/Al n-electrode

Al/Ti/Au

Al/Ti/Au

40
20

ITO/DM (18 holes)

10µm2

0
250

41 µm

ITO/DM (5 holes) Size:10

ITO/DM (1 hole)

60

300

40 µm

350

400

450

500

wavelength[nm]Ti/Al n-electrodeFig. 3. (Color online) Wavelength dispersion of the reﬂectance of the
SiO2 /AlN DM ﬁlm. In this ﬁgure, we have also included the simulation
results calculated by the matrix method.

Al/Ti/Au

DM

DM

DM

ITO(100nm)
p-GaN
(c)

ITO(100nm)
p-GaN
(d)

ITO(100nm)
p-GaN
(e)

20
Voltage[V]Fig. 2. (Color online) Top views and cross sectional views of eachelectrode structure: (a) Ni/Au, (b) ITO/Al, (c) ITO/DM (1 hole), (d) ITO/
DM (5 holes), and (e) ITO/DM (18 holes). Ti/Au pad electrodes wereformed on all electrodes. Black lines are auxiliary conductors.

16

1 holl5 holls
18 holls

1284of Ni/Au and ITO/Al electrode have been reported inRef. 15. In type III [Figs. 2(c)–2(e)], a combination of ITO(100 nm) and 10 pairs of SiO2 (60 nm)/AlN (41 nm) DM
with many holes, which were formed by a lift-oﬀ method,
as shown in Figs. 2(c)–2(e), was deposited on p-GaN by
sputtering. By forming holes and ﬁlling them with metal,
current injection can be achieved. In this study, we used a
100-nm-thick ITO ﬁlm, because the transmittance of thislayer is higher than 90% at 350 nm and its sheet resistance is
as low as 300 /. In order to increase the reﬂectance of thehole portions, a 200-nm-thick Al metal layer was depositedin the hole. Then, a Ti (10 nm)/Au (240 nm) bilayer was
deposited by electron beam evaporation for current injection.
The size of the p-electrode was 500  140 m2 . The numberof holes in SiO2 /AlN DM was varied at 1, 5, and 18, and thesize was the ﬁxed at 10  10 m2 . Ratios of the area occupiedby the holes in the p-electrode are approximately 0.02, 0.08,
and 2.5%. In addition, Ti (10 nm)/Au (240 nm) pad electrodes were formed on all electrodes for implementation.
The voltage–injection current (V –I) and light output
power–injection current (L–I) characteristics of the UVLEDs were measured using a semiconductor parameter
analyzer (Agilent 4156C). The light output powers and
EQEs of each device were measured by the power meter andintegrating sphere. Moreover, IQEs, which include current
injection eﬃciency in the active layer and LEEs of eachsample were determined by the Shockley Read Hall model
method.17–20)
3. Results and Discussion

Figure 3 shows the reﬂectance spectrum of the SiO2 /AlN
DM ﬁlm. In this ﬁgure, we have also included the simulation
results. From this ﬁgure, the reﬂectance of the SiO2 /AlN
DM ﬁlm at 350 nm was found to be approximately 98.5%.

0

0

20

40
60
Current[mA]

80

100

Fig. 4. (Color online) V –I characteristic of UV-LEDs with ITO/DM
p-electrodes with three diﬀerent numbers of holes.

Although there is a slight deviation between the simulationand experimental results, the reﬂectance of the SiO2 /AlN
DM ﬁlm is found to be higher than that of Al.
Figure 4 shows the V –I characteristics of UV LEDs
with ITO/DM p-electrodes with three diﬀerent numbers ofholes. With 18 holes, satisfactory current injection can be
conﬁrmed. Therefore, we have compared the UV LEDs with
Ni/Au, ITO/Al, and ITO/DM 18-hole electrodes in thisstudy.
A simple calculation based on the reﬂectivity results
indicates that the light output powers for the ITO/DM
electrodes are approximately 1.1 and 1.5 times higher thanthose for the ITO/Al and Ni/Au electrode assuming noabsorption, respectively. Figure 5 shows the V –I and L–I
characteristics of the 350-nm-emission UV LEDs nm withthe three diﬀerent p-electrodes. Table I also was summarized the light output powers and operating voltages of eachsamples at operation currents of 40 and 80 mA. The light
output power of the UV LED with the ITO/DM p-electrodewas 14 mW at 100 mA. The L–I characteristic shows thatthe light output power of UV LEDs with the ITO/DM
electrode was 1.17 and 1.6 times higher than those of UV
LEDs with ITO/Al and Ni/Au electrodes, respectively.
These improvements are caused by the increase in theamount of reﬂected light at the p-electrode. However, theoperating voltage at 100 mA of the UV LED with the SiO2 /08JG07-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JG07

T. Nakashima et al.

8
14

RT
DC

Voltage[V]Light Output[mW]

12

6

10
8

4

6
4
Ni/Au
ITO/AlSiO2/AlN 2

2

0

0

20

40

60

0
100

80

ﬁlled with metal, exhibited high reﬂectivity in the UV
region, good contact characteristics, and current spread,
simultaneously. The light output power of the UV LED with
the ITO/DM electrode is 1.2 and 1.6 times higher than thoseof the ITO/Al and Ni/Au electrodes, respectively.
Acknowledgements

This study was partially supported by the Program for theStrategic Research Foundation at Private Universities, 2012–2016, supported by the Ministry of Education, Culture,
Sports, Science and Technology.

Current[mA]
Fig. 5. (Color online) V –I and L–I characteristics of UV LED with ITO/
DM, ITO/Al, and Ni/Au p-electrodes.

Table I. Characteristics of the UV-LEDs with three diﬀerent p-electrodes.

Operating current = 40 mA

Operating current = 80 mA

p-electrode

Voltage(V)

Light output(mW)

Voltage(V)

Light output
(mW)

Ni/Au

3.8

2.8

4.2

6.7

ITO/Al

4.1

3.8

4.5

9.2

ITO/DM

4.5

5.0

5.1

11.1

1) Web [http://www.iwasaki.co.jp/chishiki/uv/01.html].
2) M. Kim, T. Fujita, S. Fukahori, T. Inazu, C. Pernot, Y. Nagasawa, A.

3)

4)

5)

6)
7)

Table II. EQEs, IQEs including IE, and LEEs of the UV-LEDs with three
diﬀerent p-electrodes.

p-electrode

EQE
(%)

IQEaÞ
(%)

LEE
(%)

WPE
(%)

Ni/Au

2.3

28

8

1.9

ITO/Al

2.9

29

10

2.7

ITO/DM

3.4

29

11

2.6

8)
9)
10)
11)

a) including IE.

12)
13)

AlN dielectric multilayer p-electrode was 1.1 V higher
than that of the UV LED with the Ni/Au p-electrode. The
increase in the operating voltage can be reduced by
optimizing the number and size of holes in SiO2 /AlN DM.
Table II shows the EQE, IQE including injection eﬃciency(IE), LEE, and wall plug eﬃciency (WPE) of each LED. The
EQE of LED with the ITO/DM p-electrode was 3.4%. Thisvalue was the highest among the three LEDs. Since the IQEsof the three LEDs are almost the same, we concluded that
the EQE of the UV LED with the ITO/DM p-electrode is
improved by increasing LEE.

14)
15)

16)

17)
18)

19)

4. Summary

We have investigated a novel ITO/DM reﬂective pelectrode for UV LEDs. The ITO/DM electrode with holes20)

08JG07-3

Hirano, M. Ippommatsu, M. Iwaya, T. Takeuchi, S. Kamiyama, M.
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M. Shatalov, W. Sun, A. Lunev, X. Hu, A. Dobrinsky, Y. Bilenko, J. Yang,
M. Shur, R. Gaska, C. Moe, G. Garrett, and M. Wraback: Appl. Phys.
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W. H. Sun, J. P. Zhang, V. Adivarahan A. Chitnis, M. Shatalov, S. Wu, V.
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# 2013 The Japan Society of Applied Physics

solidistatuspss

physicaPhys. Status Solidi C 11, No. 3–4, 836– 839 (2014) / DOI 10.1002/pssc.201300517

cwww.pss-c.com

current topics in solid state physics

Improvement of light extraction
efficiency of 350-nm emission
UV light-emitting diodes

Tsubasa Nakashima*,1, Kenichiro Takeda1, Motoaki Iwaya1, Tetsuya Takeuchi1, Satoshi Kamiyama1,
Isamu Akasaki1,2, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, 468-8502 Nagoya, Japan
Akasaki Research Center, Nagoya University, 464-8062 Nagoya, Japan
3
Graduate School of Engineering, Nagoya University, 464-8062 Nagoya, Japan
2

Received 19 September 2013, revised 6 December 2013, accepted 16 January 2014
Published online 12 March 2014
Keywords UV-LED, AlGaN, ITO, LLO* Corresponding author: e-mail 123434034@ccalumni.meijo-u.ac.jp, Phone: +81-52-832-1151, Fax: +81-52-832-1244

In this study, we have aimed to improve the light extraction efficiency by using p- and n- high-reflectivity indium tin oxide (ITO)/Al electrodes with a reflectivity of~77% at 350 nm, laser lift-off (LLO) method, and silicone sealing. By a combination of these three methods,

the light output power and operating voltage of UV LEDs
were improved to 23.8 mW and 6.6 V at 100 mA. Furthermore, the external quantum efficiency and light extraction efficiency were improved to 6.6 and 25.3% at
100 mA.

© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction UV light is widely used in the fieldsof medicine, biochemistry, and photolithography and insemiconductor device fabrication [1]. In particular, AlGaN-based UV-light-emitting diodes (UV LEDs) afford
many advantages such as compactness, high efficiency,
and wavelength selectivity, and therefore, these devices are
expected to replace conventional UV-light sources basedon glass lamps. Recently, many studies have reported newtechniques to achieve higher-performance AlGaN-based
UV LEDs [2-10]. However, the external quantum efficiency (EQE) of UV LEDs remains lower than those ofvisible LEDs.
Light extraction efficiency (LEE) is one of the mostimportant factors for the improvement of EQE in UV
LEDs, because the internal quantum efficiency (IQE) of
the active layer in UV LEDs has already been improved asmuch as is possible [11]. The LEE of flip-chip visible
LEDs can be improved by using high-reflectivity electrodes based on Ag [12] or Rh [13] on the back surface of
the LEDs. The LEE can also be improved by the laser liftoff (LLO) method [14] and by sealing using resin or silicone [15]. By a combination of these methods, the LEE ofvisible LEDs was improved to ~80% [16]. However, theseapproaches cannot be applied to UV LEDs.

Thus far, we have reported p- high-reflectivity electrodes made of indium tin oxide (ITO) and Al, which ishighly reflective in the UV region [17]. We have also reported the application of LLO to UV LEDs [18]. However,no study has reported on the combination of these methods
to improve the device performance of UV LEDs. In addition, UV LEDs sealed by resin or silicone have not yet
been reported.
In this study, we propose the application of a combination of p- and n-type ITO/Al reflective electrodes, LLO,
and silicone sealing for improving the device performanceof UV LEDs. Toward this end, we fabricated and characterized 350-nm UV LEDs.
2 Experiments
UV LEDs with an emission peak wavelength of 350
nm were epitaxially grown on sapphire substrates by metalorganic vapor phase epitaxy. The samples of lowdislocation-density Al0.25Ga0.75N grown on grooved
Al0.5Ga0.5N/AlN/sapphire comprised a 2.8-μm-thick nAl0.2Ga0.8N cladding layer, three pairs of unintentionally
doped GaN (3 nm), Si-doped Al0.1Ga0.9N (8 nm) multiple
quantum wells (MQWs) with a 4-nm-thick unintentionally
doped Al0.1Ga0.9N last barrier, a 20-nm-thick p-Al0.4Ga0.6N© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Contributed
ArticlePhys. Status Solidi C 11, No. 3–4 (2014)837

electron blocking layer, a 50-nm-thick p-Al0.2Ga0.8N cladding layer, and a 50-nm-thick p-GaN contact layer. The
samples were annealed at 800 °C for 5 min in N2 ambientfor p-type activation. A contact with the n-AlGaN layer
exposed by reactive ion etching was deposited by electron
beam evaporation.
We fabricated and characterized five types of flip-chip
LEDs with different device structures. Samples A–C were
fabricated to understand the usefulness of a highreflectivity electrode. Samples D and E were fabricated toinvestigate the usefulness of LLO and resin sealing. The
details of the structure of each sample are as follows:
1. Sample A:
This UV LED has a conventional Ni (10 nm)/Au (40 nm)
p-electrode that was annealed at 550 °C for 3 min in O2
ambient and a Ti (30 nm)/Al (100 nm)/Ti/Au n-electrode
that was annealed at 850 °C for 30 s in N2 ambient.
2. Sample B:
This UV LED has a high-reflectivity ITO (10 nm)/Al (150nm) p-electrode of which only ITO was annealed at 600 °C
for 5 min in N2 ambient and a Ti (30 nm)/Al (100 nm) nelectrode that was annealed at 850 °C for 30 s in N2 ambient.
3. Sample C:
This UV LED has both high-reflectivity ITO (10nm)/Al (150 nm) p- and n-electrodes that were annealed at
600 °C for 5 min in N2 ambient.
Samples E and D were used to examine the effect of
LLO and silicone sealing. We processed sample C for use

Sample A

Sample B

Sample C

Ni/Au

ITO/AlITO/Alp-GaN
p-AlGaN
MQW

p-GaN
p-AlGaN
MQW

p-GaN
p-AlGaN
MQW

Ti/Al

Ti/Al

n-AlGaNn-AlGaNn-AlGaNsapphire

sapphire

sapphire

UV light

UV light

UV light

Sample D
ITO/AlITO/AlSample E
UV light

p-GaN
p-AlGaN
MQW
ITO/Aln-AlGaN
UV light
Figure 1 Five types of flip-chip LEDs.

www.pss-c.com

in these investigations. Sample D was subjected to LLO by
irradiating a 193-nm ArF laser (pulse width: 1.5 ns, energy: 1.5 J/cm2) one time. Next, KOH wet etching was performed at 100 °C for 1 min. Figure 2 shows the scanningelectron microscopy (SEM) image of the N-face of GaN
after LLO and KOH wet etching. It is observed that the
KOH-etched N-face GaN surfaces show hexagonal conelike features with a size and height of 100–1,000 nm and
500–1,000 nm, respectively. In addition, sample E wassubjected to silicone sealing at the sample D. The silicone
sealing was of the chip-on-board type. The silicone sealing
was processed into a semi-spherical form with a diameterof 1 mm. It involved heat treatment at 150 °C for 30 min.
The voltage-injection current (V-I) and light output
power-injection current (L-I) characteristics of the LEDs
were measured using a semiconductor parameter analyzer(Agilent 4156C). The light output powers and wall plug efficiencies (WPEs) of each device were respectively measured using a power meter and an integrating sphere. Then,the IQEs, including current injection efficiency (IE) in theactive layer, and the LEEs of each sample were determinedby the Shockley–Read–Hall model method [18-21].

Figure 2 Scanning electron microscopy image of AlN
after KOH wet etching.

3 Results and discussion Figure 3 shows the V-I
and L-I characteristics of samples A–C. Table 1 summarizes the light output powers and operating voltages ofeach sample at operation currents of 50 and 100 mA. From
the L-I characteristic, the light output power of sample C
was 1.2 and 1.3 times higher than that of samples A and B,
respectively. This improvement is caused by the increasein the amount of reflected light at the p-type electrode.
Because GaN has a large absorption coefficient of ~1
× 105 cm-1 at 350 nm, approximately 60% of the reflected
light is absorbed in the p-GaN layer. Nevertheless, we confirmed that the high reflectivity electrode is effective
enough to increase the LEE. Figure 4 shows the reflectance of each electrode. Considering this absorption andthese electrode reflectances, the calculated results for samples B and C were approximately 1.3 and 1.4 times higher
than that for sample A. These calculations are in good
agreement with the experimental result. However, the op© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss

c

838

T. Nakashima et al.: Improvement of light extraction efficiency of 350-nm emission UV LEDs
100

10

12

Voltage[V]

8

10
8

6

6
4

RT, DC

0

0

20

40
60
Current[mA]

80

Reflectance[%]

40
20
0
260

280

300

320

340

360

380

Wavelength[nm]400Figure 4 The reflectance of Al, ITO/Al and Ni/Au electrode.

Sample E
Sample D
Sample C

8

20

6

15

4

10
5

2
RT, DC

0

0

20

40
60
Current[mA]

80

0
100

Table 1 Characteristics of flip-chip UV LEDs.
Operating Current = 50mA

Operating Current = 100mA

Light output[mW]

Voltage[V]

Light output[mW]

Sample A

4.4

4.2

4.9

8.7

2

Sample B

4.7

4.8

5.1

10.3

0
100

Sample C

6.1

5.5

6.6

11.4

Sample D

6.1

8.6

6.6

17.6

Sample E

6.1

11.5

6.6

23.8

Figure 3 V-I and L-I characteristics of flip-chip UV LEDs of
samples A, B, and C.

4 Summary In this study, we demonstrated the use ofa reflective electrode, LLO, and silicone sealing to improve the LEE of UV LEDs. The EQE and LEE of a 350nm-emission UV LED were improved to 6.6 and 25.3% by
a combination of both p- and n- high-reflectivity electrodes,
LLO, and silicone sealing.
Acknowledgements This study was partially supported by
the Program for the Strategic Research Foundation at Private
Universities, 2012–2016, supported by the Ministry of Education,
Culture, Sports, Science and Technology (MEXT) and a MEXT
Grant-in-Aid for Specially Promoted Research #25000011.

Table 2 EQEs, IQEs including IE, and LEEs of flip-chip UV
LEDs.

EQE
[%]

IQEa)
[%]

LEE
[%]

WPE
[%]

Sample A

1.9

26

7.6

1.7

Sample B

2.6

26

10.0

2.1

Sample C

3.1

26

11.9

1.7

Sample D

4.9

26

18.8

2.7

Sample E

6.6

26

25.3

3.6

a) Including IE© 2014 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Light Output[mW]2510Voltage
[V]

4

2

Al
ITO/AlNi/Au
Ti/Al/Ti/Au

60

Figure 5 V-I and L-I characteristics of flip-chip UV LEDs of
samples C, D, and E.
Light Output[mW]Sample C
Sample B
Sample A

80

Voltage[V]erating voltage of the UV LED with a high-reflectivity nelectrode at 100 mA was 1.6 V higher than that of the UV
LED with the Ti/Al n-electrode. This is because the ITO
and n-AlGaN formed a Schottky contact. Although the optimization of the high-reflectivity n-electrode is essential,
we found that a highly reflective n-type electrode is also
useful for improving the light output power.
Figure 5 shows the V-I and L-I characteristics of samples C–E. Because the V-I curves of each LED are almost
the same, it is suggested that there is no damage on each
LED by LLO and silicone sealing. In contrast, the light
output power of samples D and E was 1.5 and 2.0 times
higher than that of sample C, respectively. These results
suggest that LLO and silicone sealing effectively improve
the LEE in the 350-nm wavelength region. LLO increased
the light output power by more than 2.0 times in visible
LEDs [14]. However, in this experiment, the increase wasby 1.5 times, because the absorption of p-GaN occurs bymultiple reflection. We thought that this decrease in therate of increase of the light output power was affected by
the decrease in the multiple reflection effect owing to theabsorption in p-type GaN. Therefore, further improvementof the LEE can be expected if the p-type GaN layer is removed. Table 2 summarizes the EQE, IQE including IE,
LEE, and WPE of each device. The EQE and LEE of the
UV LED of sample E was 6.6 and 25.3% at 100 mA.

www.pss-c.com

Contributed
ArticlePhys. Status Solidi C 11, No. 3–4 (2014)

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GaN barrier layer dependence of critical thickness in GaInN/GaN superlattice on GaN
characterized by in situ X-ray diffraction

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REGULAR PAPER

Japanese Journal of Applied Physics 55, 05FD11 (2016)
http://doi.org/10.7567/JJAP.55.05FD11

GaN barrier layer dependence of critical thickness in GaInN/GaN superlatticeon GaN characterized by in situ X-ray diffraction
Junya Osumi1*, Koji Ishihara1, Taiji Yamamoto1, Motoaki Iwaya1, Tetsuya Takeuchi1,
Satoshi Kamiyama1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan2*E-mail: 153434005@ccalumni.meijo-u.ac.jp
Received December 29, 2015; revised February 18, 2016; accepted February 23, 2016; published online April 22, 2016
We investigated the critical thickness dependence on the GaN barrier layer thickness in a GaInN/GaN superlattice (SL). The characterization was
done by combining an in situ X-ray diffraction (XRD) system attached to a metalorganic vapor phase epitaxy rector and ex situ analyses such as
scanning electron microscopy and transmission electron microscopy. The critical thickness required for the introduction of a + c-type misﬁt
dislocations (MDs) in the GaInN/GaN SL was determined by analyzing the full width at half maximum of the in situ XRD spectrum from a GaInN/GaN SL as a function of SL periods, and we successfully found the critical thicknesses of speciﬁc different SLs.
© 2016 The Japan Society of Applied Physics

1.

Introduction

GaInN ternary alloys have bandgaps ranging from 0.65 to
3.43 eV, and are therefore used as active layers in violet,
blue, and green light-emitting diodes (LEDs).1–3) In addition,these materials can theoretically be used in the design ofa multijunction solar cell that covers the infrared to theultraviolet regions of the solar spectrum.4–10) In general, most
GaInN-based solar cells are fabricated on a GaN templateor a freestanding GaN substrate because the growth of a
thick GaInN layer is quite diﬃcult. Many previous studies
have revealed the mechanism by which defects form via
strain relaxation in GaInN=GaN heterostructures11–15) inorder to minimize defect formation and improve device
performance. We have previously reported that the superlattice (SL) is useful for improving the crystallinity of the
GaInN=GaN SL active layer.9,10,16–19) The increases in the SL
periods and InN molar fraction in the GaInN=GaN SL active
layer are required for realizing high-conversion-eﬃciency
GaInN-based solar cells. However, the crystallinity of the
GaInN=GaN SL deteriorates rapidly upon the introductionof misﬁt dislocations (MDs) and growth pits if the layer
thickness exceeds the critical thickness.4,12,13) Thus, thedetermination of the structure and growth conditions thatsuppress the generated MDs in a GaInN=GaN SL isimportant. Therefore, several experiments are needed to
explore the optimal structures unless in situ observations are
employed.
In situ monitoring in metal–organic vapor phase epitaxy(MOVPE) is the key process in device manufacturing. In
fact, the information given by this monitoring can be fed backinto the growth condition setup, and therefore, helps us tounderstand the growth mechanism. We have already reported
on the novel X-ray in situ measurements in Refs. 20 and 21.
In situ monitoring using X-rays yields a measure of thelattice constant, the nanostructure of the thin ﬁlm, and the
crystalline quality, for example. So far, we have reported thatthe use of this in situ X-ray diﬀraction (XRD). It makes
possible to accurately follow the relaxation process that
occurs in a GaInN=GaN heterostructure by analyzing the full
width at half maximum (FWHM) of the spectrum of the
GaInN layers. Moreover, we have reported the determinationof a reliable critical thickness in GaInN=GaN heterostructures
to obtain MD-free growth.22)
In the present study, we analyzed a GaInN=GaN SL using
an in situ XRD system attached to the MOVPE reactor, andthen, compared the results with those of ex situ characterization techniques such as scanning electron microscopy(SEM) and transmission electron microscopy (TEM). Thus,
we investigated the critical thickness of MDs introduced in to
the GaInN=GaN SL on GaN systems as a function of GaN
barrier layer thickness. The results indicate that the critical
thickness required for the introduction of MDs in the GaInN=GaN SL depends on the GaN barrier layer for the GaInN=GaN SL on GaN.
2.

Experimental methods

All samples were grown on c-plane sapphire and freestanding
GaN substrates using MOVPE with a horizontal face-down
2-in. × 3-wafer reactor system (Taiyo Nippon Sanso EMC
GRC-230X). Trimethylgallium (TMGa), triethylgallium(TEGa), trimethylindium (TMIn), and ammonia (NH3) wereused as Ga, In, and N source materials for the growth of GaN
and GaInN, respectively. Hydrogen (H2) and nitrogen (N2)were used as the carrier gases. The rotational speed of thesubstrate was 5 rpm. For the sapphire substrate (∼430 µm),
the GaInN=GaN SL samples were fabricated on an ∼3-µmthick GaN template grown using a low-temperature buﬀer
layer.23) For the GaN substrate (∼330 µm), the GaInN=GaN
SL samples were fabricated after the growth of an ∼1-µmthick homoepitaxial GaN layer. The threading dislocation(TD) densities of the GaN template and freestanding GaN
substrate were ∼3 × 108 cm−2 and less than 3 × 106 cm−2,
respectively. After the growth of the homoepitaxial GaN
layer at 1,050 °C on the GaN template or substrate using H2
as the carrier gas at 933 hPa, each sample was cooled to
720–750 °C and the carrier gas was changed to N2. Then,a Ga1−xInxN (2 nm)=GaN (y nm) SL was grown on the GaN
template and GaN substrate. The InN mole fraction in the
GaInN layer was varied from 0.20 to 0.25 by adjusting the
growth temperature, and the GaN barrier thickness in the
GaN layer was varied from 3 to 11 nm by adjusting the
growth time. The thickness of the GaInN well layer was set to
2 nm for all samples. We observed the GaInN=GaN SL with

05FD11-1© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD11 (2016)

J. Osumi et al.

symmetric (0002) Bragg diﬀraction using the in situ XRD
system attached to the MOVPE. The MOVPE apparatus was
equipped with two beryllium (Be) windows, which served
as viewports for the passage of X-rays focused on the sample
surface using a Johansson curved crystal mirror. By using
this method, the incidence angle of the X-rays can bechanged independently from the X-ray source. The X-rays
diﬀracted on the sample surface were detected by a onedimensional charge-coupled device (1D CCD) without
moving the substrate or the detector. This system executed
the equivalent of a (0002) 2θ=ω scan in one second during
the sample rotation without requiring the use of an analyzer
crystal. Furthermore, the 2θ value was calibrated using on
the GaN peak. Although the resolution slightly decreased
compared with that of a conventional XRD system, theequivalent of a (0002) 2θ=ω scan at a resolution of 1 arcsec
was possible. Therefore, in this system, both the tilt component and distribution of the lattice constant c can becharacterized from the FWHM. Furthermore, FWHMs of thein situ XRD spectrum were analyzed by a Gaussian ﬁtting
method. We observed the GaInN surface structure anddefects by SEM and TEM, respectively.
3.

(a)

(b)

Results and discussion

The in situ XRD spectra of Ga0.80In0.20N (2 nm)=GaN (3 nm)
SL are shown in Fig. 1(a). The FWHM obtained from thein situ XRD signal for the Ga0.80In0.20N (2 nm)=GaN (3 nm)
SL on a GaN template in the zero-order satellite peak regionis shown in Fig. 1(b). The FWHM was determined by ﬁtting
the pseudo-Voigt function for the in situ XRD spectrum.
The FWHMs for the Ga0.80In0.20N (2 nm)=GaN (3 nm) SL
decreased with increasing number of periods of SL up to 25
periods. The decrease in the FWHM is attributed to theincrease in the intensity of the satellite peak with anincreasing diﬀraction surface for the X-rays through theperiods of the GaInN=GaN SL. Hereafter, the minimum
value of FWHM was observed and the FWHM increased
after this minimum value. However, it is observed that after
the minimum point, crystal defects appeared in the GaInN=GaN SL. To investigate this phenomenon, Ga0.80In0.20N(2 nm)=GaN (3 nm) SL samples on GaN templates with
diﬀerent periods were investigated by ex situ characterization

Fig. 1. (Color online) (a) In situ XRD spectra of Ga0.80In0.20N (2 nm)=
GaN (3 nm) SL. (b) Typical FWHMs from the Ga0.80In0.20N=GaN SL on the
GaN template as a function of number of SL periods.

techniques. Plan-view SEM and TEM images for theGa0.80In0.20N (2 nm)=GaN (3 nm) SL samples on GaN templates are shown in Fig. 2. The number of periods ranges
from 10 to 40. The plan-view SEM images show signiﬁcantdiﬀerences in the size of v-pits with increasing GaInN=GaN
SL periods. The v-pit increases in size for additional periods,
as shown in Figs. 2(a)–2(c), and is homogenous for eachsample. However, the v-pits are no longer homogenous andtheir size increases with higher number of periods, as shownin Figs. 2(d) and 2(e). Note that those smaller than 200 nm

Fig. 2. (Color online) Plan-view SEM images and plan-view TEM images for Ga0.80In0.20N=GaN SL samples of diﬀerent periods on GaN templates. The
v-pit density is shown below the SEM images.

05FD11-2© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD11 (2016)

J. Osumi et al.

Fig. 3. (Color online) GaN barrier layer thickness as a function of thenumber of periods required for the generation of a + c-type MDs for diﬀerent
InN molar fractions.

appear as v-pits after the minimum point. It is reported
elsewhere that the v-pits are created by the TDs originatingfrom the underlying GaN template.13) The pit density ofeach sample is almost comparable to the TD density.
However, after 30 periods, as shown in Figs. 2(d) and 2(e),pits with diﬀerent diameters and trenches are observed. It issuggested that the larger pits originate in deeper interfaces.
The a + c-type MDs were observed after 25 periods, as
shown in Fig. 2(g), in contrast to the cases with fewer than
25 periods. Therefore, in situ XRD measures the critical
thickness for the introduction of a + c-type MDs in the
GaInN=GaN SL on the GaN system through only one-time
crystal growth.
The number of periods required for the generation of
a + c-type MDs with diﬀerent InN molar fractions in GaInN
layers as a function of thickness in the GaN barrier layer are
shown in Fig. 3. For the InN molar fraction of 0.25, thenumber of periods required for the generation of a + c-type
MDs increases from 12 to 15 on changing from the 3- tothe 7-nm-thick GaN barrier layer. Also, for the InN molar
fraction of 0.20, the number of periods required for thegeneration of a + c-type MDs increases from 25 to 29 onchanging from the 3- to the 11-nm-thick GaN barrier layer. In
contrast, it did not change in the case of 3- and 7-nm-thick
GaN barrier layers. In general, the critical thickness in the SL
increases with increasing barrier layer thickness because of
the reduction of compression stress in the GaInN well layer
upon increasing the GaN barrier layer thickness. Bykhovski
et al. reported the calculation of stress in the SL structure. In
this report, the ratio of barrier to well layer thickness islargely dependent on the critical thickness.24) Therefore, it isconsidered to be intended to suggest these results. Also, inthe case of an InN molar fraction of 0.20 with 3 and 7 nm
thicknesses of the GaN barrier layer, since the strain in
Ga0.80In0.20N is less than that in Ga0.85In0.25N, it is consideredthat the eﬀect of the GaN barrier layer is weak on a thin ﬁlm.
Moreover, on the basis of the in situ XRD results, a reliablecritical thickness for obtaining a + c-type MD-free growth
was determined.
Because we compared with a GaInN single layer on GaN,
the critical thickness of the sum of GaInN and GaN as afunction of average InN molar fraction of GaInN=GaN SL isshown in Fig. 4. We also plot the data of the critical thickness

Fig. 4. (Color online) Critical thickness of the sum of GaInN and GaN as
a function of average InN molar fraction of GaInN=GaN SL. The data for the
GaInN single layer on GaN template and GaN substrate of critical thickness
are taken from Ref. 20. The results of theoretical calculations performed
using data in Refs. 25–28 are also shown.

of the GaInN single layer on the GaN template and GaN
substrate from Ref. 20. The critical thickness required for the
introduction of a + c-type MDs in the GaInN single layer
depended strongly on the TD density of the underlying GaN
layer. In contrast, for the GaInN=GaN SL on the GaN
template and GaN substrate with the same InN molar
fraction, there is no dependence of the critical thickness on
the TD density of the underlying GaN layer. We compared
our results with theoretical calculations for GaInN=GaN
heterostructures and the critical thickness required for the
introduction of a + c-type MDs in the GaInN=GaN SL on the
GaN template and GaN substrate. These theoretical calculations, as obtained from the data in Refs. 25–28, are also
included in Fig. 4. The critical thickness of the SL behavesin accordance with the theoretical formula in Ref. 28. The
theory presented in this reference considers the interaction
between dislocations, in contrast to other theories. In the caseof the underlying GaN layer with a high TD density, it isbelieved that the critical thickness increases with the bendingof the TD. Therefore, it is considered that the critical
thickness is aﬀected by the diﬀerence in the TD density of theunderlying GaN layer. In contrast to the GaInN=GaN SL on
GaN, the critical thickness of a + c-type MDs did not changewith the TD density in the underlying GaN layer. It is thoughtthat because of the greater stress reduction eﬀect of thebarrier layer than the stress relaxation due to the bending of
the TD.
4.

Conclusions

We investigated GaInN=GaN SLs with diﬀerent GaN barrier
layer thicknesses by both in situ XRD and ex situ characterization techniques. It was demonstrated that the critical
thickness required for the introduction of a + c-type MDs canbe obtained via in situ XRD analysis. In addition, the critical
thickness required for the introduction of a + c-type MDs
in the GaInN=GaN SL was found to depend strongly on the05FD11-3© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD11 (2016)

J. Osumi et al.

GaN barrier layer thickness. For the GaInN=GaN SL on the
GaN template and GaN substrate with the same InN molar
fraction, there is no dependence of the critical thickness on
the TD density of the underlying GaN layer.
Acknowledgments

This study was partially supported by the Program for
the Strategic Research Foundation at Private Universities,
2012–2016, supported by the Ministry of Education, Culture,
Sports, Science and Technology (MEXT), the MEXT Grantin-Aid for Specially Promoted Research No. 25000011, andthe MEXT Grant-in-Aid for Scientiﬁc Research A No.
15H02019.

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Extremely Low-Resistivity and High-Carrier-Concentration Si-Doped Al0.05Ga0.95N

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Applied Physics Express 6 (2013) 121002
http://dx.doi.org/10.7567/APEX.6.121002

Extremely Low-Resistivity and High-Carrier-Concentration Si-Doped Al0:05 Ga0:95 N
Toru Sugiyama1 , Daisuke Iida1 , Toshiki Yasuda1 , Motoaki Iwaya1 ,Tetsuya Takeuchi1 , Satoshi Kamiyama1 , and Isamu Akasaki1;2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
E-mail: iwaya@meijo-u.ac.jp
2

Received August 29, 2013; accepted November 7, 2013; published online November 27, 2013
We discovered that Si-doped Al0:05 Ga0:95 N can be used to realize an extremely low-resistivity n-type layer at room temperature. In Si-doped GaN,
a resistivity of 2:7  10 3  cm with a carrier concentration of 4:0  1019 cm 3 was almost saturated. In contrast, Si-doped Al0:05 Ga0:95 N with aminimum resistivity of 5:9  10 4  cm was produced with an electron concentration and electron mobility of 1:4  1020 cm 3 and 70 cm2 /V s,
respectively. We confirmed a reduction in the differential resistance of a violet light-emitting diode with a high external quantum efficiency by using
this Si-doped Al0:05 Ga0:95 N. # 2013 The Japan Society of Applied Physics

roup-III nitride semiconductors are used as lightemitting diodes (LEDs)1–4) and laser diodes5,6) (LDs)having a wide wavelength range. High-brightness
GaInN-based visible LEDs are presently commercialized
worldwide,7–9) but from a practical viewpoint, the wall plugeﬃciency (WPE) of such devices needs improvement. The
external quantum eﬃciency (EQE) of these devices is alreadysuﬃciently high;10,11) therefore, the device resistance, which
includes the sheet, bulk, and contact resistances of n- and ptype layers, needs to be reduced to improve the WPE. In
particular, the low-resistivity (low-) n-type layer is important because commonly used nitride-based LEDs require
for current spreading layers and for simple contacting. A
number of studies have reported on the reduction of  in ntype GaN.12–16) Fritze et al.16) managed to reduce  for the Si
doping of GaN to 2:5  103  cm at room temperature(RT). They stated that a  and free electron concentration (ne )of less than 2:5  103  cm and higher than 1:9  1019
cm3 , respectively, could not be achieved because 3D growthoccurs by excess Si-doping. In contrast to Si-doping, a highne (2:9  1020 cm3 ) and low- (6:3  104  cm) GaN
layer could be realized by Ge doping.16)
In this study, we demonstrated n-type cladding layer ofSi-doped Al0:05 Ga0:95 N with low- and high-ne with smoothsurface. We improved the WPE of a violet LED by using
this Si-doped Al0:05 Ga0:95 N.
All samples were grown by metal organic vapor-phase
epitaxy. Figure 1 shows a schematic view of the sample
structure. Trimethylindium, trimethylaluminum, trimethylgallium, triethylgallium, and ammonia were used as thesource gases. C-plane sapphire was used as a substrate;
it was covered with a low-temperature-deposited buﬀer
layer.17) An unintentionally doped 3-m-thick GaN layer
and a 1.5-m-thick Si-doped Al0:05 Ga0:95 N layer or GaN
layer were then sequentially grown. Silane (SiH4 , 50 ppm
diluted by H2 ) was used as a dopant source.18) The carrier
concentration of the Si-doped Al0:05 Ga0:95 N and GaN was
adjusted from 8:0  1018 to 1:5  1020 cm3 . We measured
the van der Pauw Hall eﬀect to determine ne , mobility (e ),and  of each sample at RT. The two-dimensional electron
gas at the heterointerface between Al0:05 Ga0:95 N and GaN
was ignored because the sheet resistance of each sample wasvery low. We also characterized the surface morphology byscanning electron microscopy (SEM).
We fabricated a violet LED with a 390-nm-emission
GaInN multiple quantum well (MQW) active layer on theG

Si-doped GaN or
Al0.05Ga 0.05N (1.5 μm)
unintentionally doped GaN(~3 μm)LT-buffer layer
Sapphire substrate
Fig. 1. Schematic view of the sample structure.

Cr/Ni/Au
Ni/Au
p-GaN
p-Al0.05Ga0.95N
p-Al0.3Ga0.7N
Ga0.92In0.08N/Al0.05Ga 0.95N
MQW

PAD electrode

Cr/Ni/Au

Si doped Al0.05Ga0.95N

p-electrode
350 μmUnintentionally doped GaN
LT-buffer layer
Sapphire substrate

Fig. 2.

n-electrode

Schematic structure and top view of the violet LED.

Si-doped Al0:05 Ga0:95 N layer. This LED structure wascomposed of a 2.0-m-thick Si-doped n-Al0:05 Ga0:95 N layer,
three pairs of unintentionally doped Ga0:92 In0:08 N (2.7 nm)/Al0:05 Ga0:95 N (12 nm) MQWs, a 20-nm-thick Mg-doped
p-Al0:3 Ga0:7 N electron-blocking layer, a 100-nm-thick
p-Al0:05 Ga0:95 N cladding layer, and a 10-nm-thick p-GaN
contact layer. Figure 2 shows the schematic device structure
and top view of this LED. After the samples were annealed
at 800  C for 10 min in ambient air for p-type activation, we
fabricated a face-up LED. A Cr/Ni/Au ohmic electrode wasdeposited by electron beam evaporation on the n-Al0:05 Ga0:95 N layer, which was exposed by inductively coupledplasma etching. A Ni/Au ohmic semitransparent electrodeand Cr/Ni/Au pad electrode were deposited on the p-GaN
layer. We prepared two samples with diﬀerent Si-doped
Al0:05 Ga0:95 N (ne ¼ 1:0  1019 and 1:6  1020 cm3 ). The

121002-1# 2013 The Japan Society of Applied Physics

Appl. Phys. Express 6 (2013) 121002

T. Sugiyama et al.

(a)

RT

(b)

-2

10
5um

Fig. 3. Plan-view SEM image of (a) Si-doped GaN(ne ¼ 3:0  1019 cm3 ) and (b) Si-doped AlGaN (ne ¼ 1:4  1020 cm3 ).

-3

10

GaN
Al0.05Ga0.95N
18

19

10

20

10
-3
ne [cm ]10

Fig. 4. ne dependence of  in Si-doped GaN and Al0:05 Ga0:95 N.

μ e [cm /Vs]

RT

2

100

rough
GaN
surfaceAl0.05Ga0.95N

10

18

19

10

20

10

10
-3

ne [cm ]ne dependence of e in Si-doped GaN and Al0:05 Ga0:95 N.

Fig. 5.

2

Current density [A/cm ]
5

0

27

54

81

108

135
25

4

20

3

15

2

10
19

0

-3

1.0 x 10 [cm ]20-3
1.6 x 10 [cm ]

1

0

20

40

60

80

5

Output power [mW]

RT, DC
Voltage [V]device size of each LED was 350  350 m2 . The voltage–
injection current (V –I) and light output power–injection
current (L–I) characteristics of the LEDs were measured
using a semiconductor parameter analyzer (Agilent 4156C).
Each LED was characterized with a bare chip. The light
output powers and WPEs of each device were measured
using the photodiode.
Figure 3 shows the plan-view SEM image of the Si-doped(a) GaN (ne ¼ 3:0  1019 cm3 ) and (b) Al0:05 Ga0:95 N (ne ¼
4:0  1020 cm3 ). Signiﬁcant diﬀerences in surface morphology were identiﬁed through these ﬁgures. In Fig. 3(a),the Si-doped GaN showed island growth (size, 1.5 m toseveral tens of micrometers). Because ne of Si-doped GaN
up to 1:0  1019 cm3 can be obtained on a smooth surface,
we concluded that the surface of this sample was roughbecause of excess Si. This result is almost in agreement withthat reported in Ref. 16. In contrast, a smooth surface wasobtained in Al0:05 Ga0:95 N despite the Si concentration being
too high; we were unable to conﬁrm the existence of cracks
in all Al0:05 Ga0:95 N samples. We concluded that a smoothsurface can be obtained by the addition of AlN even though
for a high Si concentration.
Figures 4 and 5 show  as a function of ne and e as afunction of ne , respectively, in Si-doped Al0:05 Ga0:95 N and
GaN. ne of the Si-doped GaN was unambiguously increased
up to 4:0  1019 cm3 . However,  of 2:7  103  cm withne of 4:0  1019 cm3 was close to saturation because e
rapidly decreased with deteriorating surface roughness.
These results are similar to those reported in Ref. 16. In contrast, In contrast, ne of Al0:05 Ga0:95 N continued to increase to1019 cm3 and further to 1:4  1020 cm3 . The minimum 
in Si-doped Al0:05 Ga0:95 N was 5:9  104  cm. This  islower than that of GaN with Ge doping.16) ne and e of
n-Al0:05 Ga0:95 N in the minimum- sample were 1:4  1020
cm3 and 70 cm2 /V s, respectively. Because 3D growth issuppressed using Al0:05 Ga0:95 N even under high Si concentration, e of this n-Al0:05 Ga0:95 N did not decrease signiﬁcantly. This result shows the potential of realizing a low-
n-type layer by using Al0:05 Ga0:95 N.
Figure 6 shows the measurement results of I–L and I–Vof the violet LED on diﬀerent Si-doped Al0:05 Ga0:95 N (ne :1:0  1019 and 1:6  1020 cm3 ). We conﬁrmed a signiﬁcantdiﬀerence in the I–V curve of each LED. By increasing the
Si concentration (ne in n-Al0:05 Ga0:95 N was changed from
1:0  1019 and 1:6  1020 cm3 ), we found that the operating voltage was reduced by 1 V at 100 mA current injection. Moreover, the diﬀerential resistance at 100 mA current
injection (series resistance) of each LED was reduced from
14 to 7.2 . According to a simple calculation, approxi-

roughsurfaceρ [Ω cm]

50um

0
100

Current [mA]
Fig. 6. I–V and I–L measurements of the LED on Si-doped
Al0:05 Ga0:95 N.

121002-2# 2013 The Japan Society of Applied Physics

Appl. Phys. Express 6 (2013) 121002

T. Sugiyama et al.

can be realized at RT. In contrast,  and ne in Si-dopedn-type GaN of less than 2:5  103  cm and higher
than 1:9  1019 cm3 , respectively, could not be achieved
because 3D growth occurs by excess Si-doping. Furthermore, the reduction of the series resistance of the LED witha high internal quantum eﬃciency was possible by using thisunderlying layer. Therefore, this Si-doped n-Al0:05 Ga0:95 N
underlying layer is extremely useful for the realization ofhigh-performance GaInN-based LEDs.

2

Current density [A/cm ]Wall plug efficiency [%]

8

0

27

54

81

108

135

RT, DC
6

4

2

19

-3

1.0x10 [cm ]20-3
1.6x10 [cm ]

0
0

20

40

60

80

1) H. Amano, M. Kito, K. Hiramatsu, and I. Akasaki: Jpn. J. Appl. Phys. 28

100(1989) L2112.

Current [mA]2) M. Koike, N. Koide, S. Asami, J. Umezaki, S. Nagai, S. Yamasaki, N.

Shibata, H. Amano, and I. Akasaki: Proc. SPIE 3002 (1997) 36.
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Fig. 7. WPEs for injection current.

34 (1995) L797.
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Mukai: Jpn. J. Appl. Phys. 34 (1995) L1332.

mately 4  is caused by the p-type layer (contact and bulk
resistance) among the series resistance. Because  was
reduced to approximately 1/5 times by ne increases from
1:0  1019 and 1:6  1020 cm3 , this reduction in diﬀerential
resistance can be partially explained by the decrease in the
sheet resistance of the n-Al0:05 Ga0:95 N layer. Because thelight output power of each LED for the low-current injection
was almost the same, as indicated by the I–L curves, thequality of the active layer is suﬃciently high even whenusing an underlying layer with a high Si concentration. In
contrast, the light output power in violet LEDs when using
an underlying layer with a high Si concentration at a highcurrent injection was high approximately 5%. Cause of thisdiﬀerence can be considered heat generation,21) decreaseof the current injection eﬃciency by current over ﬂow,22)
and droop.23) Among these, we considered the decreaseof injection eﬃciency and droop are irrelevant, because
structure of the p-type layer and the GaInN based active
layer in both devices is the same. Also, since we have notseen a noticeable diﬀerence in the light emission pattern ofeach LEDs, we attributed this to the fact that heat generation is reduced by reductions in diﬀerential resistance.
Therefore, this n-Al0:05 Ga0:95 N underlying layer with ahigh Si concentration is useful for a high-current-operation
LED. Figure 7 shows the WPE as a function of the injection current. The WPE of the violet LED when using an
underlying layer with a high carrier concentration wasaround 1.15 times as high as that when using an underlying
layer with a low carrier concentration. Therefore, this Sidoped n-Al0:05 Ga0:95 N underlying layer is extremely usefulfor realizing high-performance GaInN-based LEDs. In addition, because AlGaN is also useful as a UV light-emitting
device material,19,20) we expect that our results can beapplied to UV and deep-UV LEDs and laser diodes.
In conclusion, we discovered that Si-doped n-type Al0:05 Ga0:95 N with extremely low resistivity (5:9  104  cm)5) I. Akasaki, H. Amano, S. Sota, H. Sakai, T. Tanaka, and M. Koike: Jpn. J.

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7)

8)

9)

10)
11)
12)
13)
14)
15)

16)
17)
18)
19)

20)

21)

22)

23)

121002-3

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# 2013 The Japan Society of Applied Physics

Small current collapse in AlGaN/GaN
HFETs on a-plane GaN self-standing
substrate

solidistatuspss

physicaPhys. Status Solidi C 9, No. 3–4, 875–878 (2012) / DOI 10.1002/pssc.201100397

cwww.pss-c.com

current topics in solid state physics

T. Sugiyama*,1, Y. Honda1, M. Yamaguchi1, H. Amano1,3, Y. Isobe2, M. Iwaya2, T. Takeuchi2,
S. Kamiyama2, I. Akasaki2,3, M. Imade4, Y. Kitaoka4, and Y. Mori4
1

Graduate School of Engineering, Nagoya University, C3-1 Furo-cho, Chikusa-ku, Nagoya 464-8603, Japan
Faculty of Science and Technology, Meijo University, 1-501 Shiogamaguchi, Tempaku-ku, Nagoya 468-8502, Japan3
Akasaki Research Center, Nagoya University, Furo-cho, Chikusa-ku, Nagoya 464-8603, Japan
4
Graduate School of Electrical Engineering, Osaka University, Suita 565-0871, Japan
2

Received 8 July 2011, revised 20 December 2011, accepted 24 January 2012
Published online 17 February 2012
Keywords GaN, AlGaN/GaN, HFET, nonpolar, a-plane, current collapse* Corresponding author: e-mail takayu_s@nuee.nagoya-u.ac.jp

We measured drain bias stress effects and current collapse in AlGaN/ GaN heterostructure field-effect transistors (HFETs) on a-plane and c-plane GaN substrates. An
a-plane AlGaN/GaN HFET (a-HFET) shows small current collapse with a threshold voltage (Vth= -1.8 V).
On the other hand, a c-plane HFET (c-HFET) with thesame barrier thickness (20 nm) shows a small current collapse, although Vth was negatively large (Vth= -4 V).

Current collapse in a-HFET was not large compared withthat in conventional c-HFET on GaN. A c-HFET on sapphire was also measured. The current collapses in HFETs
on sapphire were larger than that in HFETs on a GaN
substrate. The current collapses in the thin-barrier cHFET (Vth = -1.8 V) were particularly large. Therefore,
an a-plane device is promising for a small or positive Vthwith small current collapse.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction AlGaN/GaN HFETs are expected forhigh-efficiency and high-power switching devices in power inverter circuits. Normally off mode operation is necessary to realize a fail-safe operation in power switching applications. Studies on normally off operation have been reported by many groups [1-5]. Nonpolar AlGaN/GaN
HFETs are also expected as a structure to realize normallyoff operation. The heterojunction on a nonpolar plane(a-plane and m-plane) is not affected by piezoelectric andspontaneous polarizations. Therefore, two dimensional
electron gases (2DEGs) channel at the hetero-interface canbe controlled merely by Si doping. However, dislocationand stacking fault density in nonpolar GaN on a highly
mismatched substrate are huge, which results in very low
mobility. Recently, HFETs on m-plane and a-plane GaN
substrates have been reported [5-7]. High drain current andsmall threshold voltage (Vth) were realized using a nonpolar HFET on GaN substrates.
The current collapse in HFETs caused by electron traps
has been reported [8, 9]. The surface states and carrier

generation mechanisms of c-plane and a-plane HFETs aredifferent. Therefore, it is meaningful to characterize thecurrent collapse property in a-plane HFETs (a-HFETs). In
this study, we measured drain bias stress effects and current collapse in HFETs on a-plane and c-plane GaN substrates and conventional HFETs on sapphire. The relationships the FET characteristics, the substrate and device
structure were measured.
2 Experiments
2.1 Device structure Figure 1 shows the device
structures. We fabricated 4 samples. All the HFETs weregrown by MOVPE. The a-HFET was grown on an a-plane
GaN liquid phase epitaxial (LPE) substrate. An unintentionally doped thick GaN layer and an Al0.3Ga0.7N barrierof 21 nm thickness were grown. The AlGaN barrier contained a Si-doped layer. As references, an Al0.3Ga0.7N(22 nm)/GaN HFET on a free-standing c-GaN substrateand conventional Al0.3Ga0.7N (10 and 23 nm)/GaN HFETs
on c-sapphire were also grown. After the growth, the sam© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

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T. Sugiyama et al.: Small current collapse in AlGaN/GaN HFETs

ple was etched for mesa isolation by Cl2-based reactive ion
etching (RIE). Ti/Al/Ti/Au was used as the source-drainelectrode, while Ni/Au was used as the schottky gate metal
for the metal-semiconductor (MES) structure. The gate
length (LG) and the distance between the source and drain(LGD) of all the devices were 2 and 8 μm, respectively. The
current collapse in all the devices was measured without
passivation.

2.2 Current collapse measurements The drain
stress effects and current collapse of HFETs were measured in a dark room. The current collapse measurements
were done by placing the substrate on a quartz plate. A parameter analyzer 4155C (Agilent) was used for all themeasurements.
Figure 2 shows the transfer characteristics (VGS-IDS) of
4 samples with VDS = 5 V. Vth and IDS of the a-HFET and
c-HFET (Al0.3Ga0.7N: 10 nm) were almost the same
(Vth=-1.8 V). On the other hand, the c-HFETs with a thick
AlGaN barrier (Al0.3Ga0.7N: 22 and 23 nm) showed high
IDS but the Vth values were negatively large.

a-HFET

Source

Gate

Drain

a-HFET
c-HFET on GaN(AlGaN:22 nm)
c-HFET on Sapp(AlGaN:10 nm)
c-HFET on Sapp(AlGaN:23 nm)

0.20

Gate

Drain

0.15
I DS (A/mm)

c-HFET on GaN
thick barrier
Source0.10

0.05
c-HFET on Sapp.
thin barrier

Source

Gate

Drain

0.00

-6

-4

-2
V GS (V)02

Figure 2 Transfer characteristics (VGS-IDS) with VDS=5 V.

c-HFET on Sapp.
thick barrier

SourceGate

Figure 1 Schematic view of HFETs.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Drain

3 Results The drain stress bias was maintained for
30 s with VGS = -8 V. The time shift of the drain current
was measured when the gate and drain were turned on(VGS = 0 V, VDS = 5 V). Figure 3 shows the time shift of thedrain current with stress bias in a-HFETs. IDS decreased
with increasing stress bias. The IDS was not quickly recovered with a very long time constant. However, the IDS wasrecovered by exposure to room light. Therefore, the originof the current collapse is not device damage but the electron traps. This current collapse is considered to originate
from deep traps as indicated by the long time constant [10].
Figure 4 shows the stress bias and stress time dependenceof the normalized IDS at 10 s after the stress bias. The IDS
and current collapses depend not on the stress time but on
the stress voltage. This trend is the same in all HFETs.
Thus, the following current collapse measurements werecarried out with the stress time of 30 s.
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Contributed
ArticlePhys. Status Solidi C 9, No. 3–4 (2012)877

switching, by GaN substrates and a thick AlGaN barrier. A
small current collapse is very important in a high-speed
switching device. The electric field was decreased by conductive GaN substrates, and the trapped electrons werealso decreased [11]. A virtual gate is generated by the coulomb force of the trapped electrons. Therefore, the effectsof current collapse in the HFET with a thin barrier layer
are large.

0.15
beforeIDS (A/mm)

0.10
Vstress=20V
30V
40V
50V

0.00
0.0

2.5

5.0

1.0
0.8

7.5

Normalized IDS

0.05

10.0

Time (s)Figure 3 Time shift of IDS in a-HFET.

0.6
0.4
a-HFET
c-HFET on GaN (AlGaN:22 nm)
c-HFET on Sapp (AlGaN:10 nm)
c-HFET on Sapp (AlGaN:23 nm)

0.2
0.0

0

10

0.7

1.0

0.6
30

20
30

60

Ti
m

40

90

e(
s)

120

50

B
ss
e
r
St

ia

V)
s(

Figure 4 Normalized IDS at 10 s after the stress bias and stresstime.

Figures 5 and 6 show the decrease in drain current ineach HFET with stress bias. Figure 5 shows the normalized
IDS at 10 s after stress off and FET turn on. The decrease in
IDS in the a-HFET is small next to that in the c-HFET on
GaN (AlGaN: 22 nm), and the same as that in the c-HFET
on sapphire (AlGaN: 23 nm). In contrast, decrease in IDS in
the c-HFET on sapphire (AlGaN: 10 nm) is very large.
Figure 6 shows the normalized IDS at 0.1 s. The differenceof the normalized IDS was large immediately after the
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50

a-HFET on GaN )c-HFET on GaN (AlGaN:22 nm)
c-HFET on Sapp. (AlGaN:10 nm)
c-HFET on Sapp. (AlGaN:23 nm)

0.8

St r
ess

40

Figure 5 Normalized IDS at 10 s for each HFET.

0.9

0.8
Normalized IDS

Normalized IDS

1.0

20
30
V DS (V)

0.6
0.4
0.2
0.0

0

10

20
30
V DS (V)40

50

Figure 6 Normalized IDS at 0.1 s for each HFET.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss

c

878

T. Sugiyama et al.: Small current collapse in AlGaN/GaN HFETs

The a-HFET with a conventional structure and without
passivation showed a small Vth and a small current collapse
compared with c-HFETs. Thus, improvements of the performance of the a-HFET are expected by passivation andwith the reform of the device structure design [12, 13].

[2] Y. Uemoto, M. Hikita, H. Ueno, H. Matsuo, H. Ishida, M.
Yanagihara, T. Ueda, T. Tanaka, and D. Ueda, IEEE Trans.
Electron Devices 54, 3393-3399 (2007).
[3] T. Mizutani, M. Ito, S. Kishimoto, and F. Nakamura, IEEE
Electron Device Lett. 28, 549-551 (2007).
[4] M. Kanamura, T. Ohki, T. Kikkawa, K. Imanishi, T. Imada,
A. Yamada, and N. Hara, IEEE Electron Device Lett. 31
4 Summary The drain bias stress effects and current
189-191 (2010).
collapse in a-plane and c-plane HFETs were measured.
[5] T. Fujiwara, S. Rajan, S. Keller, M. Higashiwaki, J. S.
The c-HFET with thick barrier thickness (22 nm) on GaN
Speck, S. P. DenBaars, and U. K. Mishra, APEX 2, 011001
shows small current collapse; however, Vth was negatively(2009).
large. On the other hand, a c-HFET with a thin AlGaN bar[6] T. Fujiwara, S. Keller, M. Higashiwaki, J. S. Speck, S. P.
rier (10 nm) on sapphire shows a small Vth because the
DenBaars, and U. K. Mishra, APEX 2, 061003 (2009).
2DEG sheet carrier density is small. However, in contrast,[7] Y. Isobe, H. Ikki, T. Sakakibara, M. Iwaya, T. Takeuchi, S.
the current collapse in the thin-barrier c-HFET was large.
Kamiyama, I. Akasaki, T. Sugiyama, H. Amano, M. Imade,Because the surface was very close to the 2DEG channel,
Y. Kitaoka, and Y. Mori, APEX 4, 064102 (2010).
[8] R. Vetury, N. Q. Zhang, S. Keller, and U. K. Mishra, IEEE
maximum electric field concentration is high and theTrans. Electron Devices 48, 560-566 (2001).
2DEG sheet carrier density is low. Therefore, the effect of[9] H. Hasegawa, T. Inagaki, S. Ootomo, and T. Hashizume, J.
trapped electrons was large.
Vac. Sci. Technol. B 21, 1844-1855 (2003).
In contrast, Vth of a-HFETs were small, because careerconcentration in a-HFET structure was low, however the [10] J. Joh and J. del Alamo, IEEE Trans. Electron Devices 58,
132-140 (2011).
effects of current collapse were small. These characteris[11] W. Saito, T. Nitta, Y. Kakiuchi, Y. Saito, K. Tsuda, I.
tics in a-HFET are contrast with that in c-HFET. Moreover,
Omura, and M. Yamaguchi, IEEE Electron Device Lett. 28,
there are differences in device structures, conductive sub523-525 (2007).
strates and Si delta-doped AlGaN barrier layer. These dif- [12] M. F. Romero, A. Jiménez, J. Miguel-Sánchez, A. F. Braña,
ferences also effected the current collapse [14].
F. González-Posada, R. Cuerdo, F. Calle, and E. Muñoz,
The a-plane AlGaN/GaN HFET on GaN substrates
IEEE Electron Device Lett. 29, 209-211 (2008).
shows a small current collapse and Vth. Current collapse in [13] N. Tsurumi, H. Ueno, T. Murata, H. Ishida, Y. Uemoto, T.
a-HFET was not bad compared with that in the convenUeda, K. Inoue, and T. Tanaka, IEEE Trans. Electron Detional c-HFET, in spite of a-HFETs were nonpolar plane
vices 57, 980-985 (2010).
growth which have a mixture surface of Ga and N. Thus, [14] G. Meneghesso, G. Verzellesi, F, Danesin, F. Rampazzo, F.
Zanon, A. Tazzoli, M. Meneghini, E. Zanoni, IEEE Trans.
the a-plane device is promising for the realization of aElectron Devices 53, 332-343 (2008).
small or positive Vth and small current collapse.

Acknowledgements
This work was supported by the New Energy and Industrial
Technology Development Organization Project, Nanotechnology
and Materials Technology, "Development of Nitride-Based Semiconductor Single Crystal Substrate and Epitaxial Growth Technology," and the second-stage Knowledge Cluster Initiative, Tokai Region Nanotechnology Manufacturing Cluster in Japan.

References[1] N. Tsuyukuchi, K. Nagamatsu, Y. Hirose, M. Iwaya, S.
Kamiyama, H. Amano, and I. Akasaki, Jpn. J. Appl. Phys.
45, 319-321 (2006).

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Trench-Shaped Defects on AlGaInN Quantum Wells Grown under Different Growth Pressures

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REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JB27
http://dx.doi.org/10.7567/JJAP.52.08JB27

Trench-Shaped Defects on AlGaInN Quantum Wells Grown
under Different Growth Pressures
Tomoyuki Suzuki1 , Mitsuru Kaga1 , Kouichi Naniwae2 , Tsukasa Kitano2 , Keisuke Hirano1 ,Tetsuya Takeuchi1 , Satoshi Kamiyama1;2 , Motoaki Iwaya1 , and Isamu Akasaki1;31Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
EL-SEED Corporation, Nagoya 468-0073, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received October 13, 2012; revised March 11, 2013; accepted March 24, 2013; published online July 22, 2013
We obtained high-quality AlGaInN/GaN quantum wells (QWs) with a 385 nm emission by adding Al sources to GaInN QWs grown at a low growthpressure of 150 mbar. When AlGaInN QWs were grown at a relatively high growth pressure of 400 mbar, a considerable amount of peculiar
trench-shaped defects were observed on the surface of the AlGaInN QWs. We found that the trench defects acted as nonradiative centers,
leading to low photoluminescence (PL) intensities. Our experiments reveal that a low growth pressure is effective in suppressing the trench defectformation, resulting in high-quality AlGaInN QWs. # 2013 The Japan Society of Applied Physics

1. Introduction

Table I. Materials of well/barrier and growth conditions at well layers.

AlGaInN quarternary alloy has an important ﬂexibility inthat the band-gap energy can be designed while the lattice
matching condition is satisﬁed. This ﬂexibility is very
useful for obtaining high-quality heterointerfaces, which arecommonly required in various semiconductor devices.1–3)
However, it is diﬃcult to grow the AlGaInN alloy since thetypical growth conditions of AlN and InN are very diﬀerent,such as growth temperature and pressure. A high growthtemperature and a low growth pressure are necessary for the
AlN growth to reduce parasitic reactions between Al sources
and N sources and enhance Al adatom migrations on thesurfaces.4–7) In contrast, a low growth temperature and ahigh growth pressure are required for the InN growth toincorporate In atoms into epitaxial layers.8–11) The two
growth parameters are completely opposite.
Recently, a novel monolithic white LED fabricated using
a combination of a ﬂuorescent SiC substrate and a nitride
semiconductor near UV-LED structure was proposed.12–15)
The UV-LED should have a peak wavelength of 385 nm toexcite carriers in the ﬂuorescent SiC substrate. To obtain
385 nm emission, Ga0:95 In0:05 N quantum wells (QWs) are
typically used.16) Unfortunately, the 0.05 InN mole fraction
is too small to make the most of the In inhomogeneity for ahigher radiative eﬃciency. The eﬀect of In inhomogeneity
has been observed in purple to blue GaInN-based LEDs inwhich the GaInN QWs contain more than 0.1 InN mole
fractions. We then proposed to use AlGaInN QWs for
385 nm emission. AlGaInN with 385 nm emission shouldhave relatively large (0:1) InN and AlN mole fractions
simultaneously. The concept is to obtain not only suﬃcient
In inhomogeneity with about 0.1 InN mole fraction but
also a large band-gap corresponding to 385 nm emission
with about 0.1 AlN mole fraction. So far, the AlGaInN
alloys with a few percent InN mole fractions have been
investigated as active regions for the emission of less than
365 nm.17–22)
In this study, AlGaInN QWs targeting high-eﬃciency
385 nm emission are fabricated and characterized. Our
approach to obtaining the AlGaInN QWs is to supply
additional Al sources to our typical GaInN QW growth
conditions. We found that peculiar trench-shaped defects

SampleMaterials ofwell/barrier
layers

A1

Ga0:9 In0:1 N/GaN

400

32

37

0

A2

AlGaInN/AlGaN

400

32

37

9

B1

AlGaInN/GaN

400

4

37

9

B2

AlGaInN/GaN

400

1

37

9

B3

Al0:07 Ga0:93 N/GaN

400

0

37

9

Pressure
TMIn
TEGa
TMAl
(mbar) (mol/min) (mol/min) (mol/min)

C1

Ga0:87 In0:13 N/GaN

150

38

46

0

C2
C3

AlGaInN/GaN
Al0:86 In0:14 N/GaN

150
150

38
38

46
0

9
9

showing a strong correlation with nonradiative recombinations were observed on the surfaces of the AlGaInN QWs.23)
A key growth parameter is discussed to suppress theformation of the trench-shaped defects.
2. Experimental Methods

A sample structure in the experiments is ﬁve QWs grown
on a Si-doped GaN/undoped GaN/sapphire template bymetalorganic vapor phase epitaxy (MOVPE). The well andbarrier thicknesses were estimated to be about 3 and 12 nm,
respectively, from X-ray diﬀraction patterns and growthtime. For the ﬁve QW structures, trimethylaluminum(TMAl), triethylgallium (TEGa), trimethylindium (TMIn),ammonia (NH3 ), and nitrogen were used as Al, Ga, In, andN sources and a carrier gas, respectively. The growthtemperature was 780  C to obtain reasonably high In
incorporation. A total of eight samples in three sets wereprepared, as listed in Table I. The ﬁrst set of two samples(A1 and A2) is for growing AlGaInN QWs on the basis of astandard GaInN QW. Sample A1 was our standard GaInN/GaN QWs grown at 400 mbar. Then, Sample A2 was grownunder the same growth condition as Sample A1 but with anadditional TMAl ﬂow at the wells and barriers, resulting in
AlGaInN/AlGaN QWs. The second set of three samples (B1
to B3) is for growing AlGaInN/GaN QWs with smaller
TMIn ﬂows to observe the TMIn ﬂow dependence of surface
morphologies. Samples B1 to B3 were grown under the

08JB27-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JB27

T. Suzuki et al.

same growth conditions as Sample A2 but with smaller
TMIn ﬂows. Actually, no TMIn ﬂow was used for
Sample B3, meaning that the nominal well material was
AlGaN. The third set of three samples (C1 to C3) is for
growing AlGaInN QWs on the basis of another standard
growth condition of GaInN QWs. Sample C1 was GaInN/GaN QWs grown under a lower growth pressure of
150 mbar. Samples C2 (AlGaInN/GaN) and C3 (AlInN/
GaN) were grown under the same low growth pressure butwith TMAl ﬂows. In Sample C3, only TMIn and TMAlﬂows were supplied to the wells.
For the ternary alloys, such as Samples A1, B3, C1, andC3, the mole fractions were determined from X-ray
diﬀraction patterns as listed in Table I. Even though thegrowth temperature and In=ðGa þ InÞ supply ratio were thesame between A1 and C1, the InN mole fractions wereslightly diﬀerent. This could be due to diﬀerences in the
total ﬂow rate as well as growth pressure.
For the quarternary alloys, such as A2, B1, B2, and C2,
the mole fractions are unknown owing to insuﬃcient
information. However, note that the InN mole fractions of
C1 (GaInN) and C3 (AlInN) are almost the same. The
growth conditions of both the samples are the same in termsof growth temperature, growth pressure, and the amount of
TMIn supply but diﬀer in the amount of other group-III (Gaor Al) source supplies. Thus, in this experiment, the InN
mole fractions of A2 and C2 (AlGaInN) may be similar to
those of A1 and C1 (GaInN), respectively.
The surface morphologies of all the samples werecharacterized by atomic force microscopy (AFM). In
addition, surface images were obtained by scanning electron
microscopy (SEM) and room-temperature panchromatic(160–930 nm) cathodoluminescence (CL) analysis in thesame areas to correlate the surface morphologies withluminescence intensities. Also, room temperature photoluminescence (PL) spectra were measured to evaluate theoptical qualities of the QW samples.
3. Results and Discussion

Room temperature PL revealed that the AlGaInN QWSample A2 showed a 385 nm peak wavelength as designed
but also a much lower PL intensity than the standard GaInNQW Sample A1. Note that AlGaInN QW Sample A2
contains AlGaN barrier layers, as listed in Table I, foran equivalent carrier conﬁnement to that of GaInN QWSample A1. Thus, the lower PL intensity of Sample A2
could suggest a poor quality of the AlGaInN QW material.
To investigate this further, surface morphology images
obtained by AFM and SEM, and CL images were observed,
as shown in Fig. 1(a) for the GaInN Sample A1 andFig. 1(b) for AlGaInN Sample A2. The SEM and CL images
were taken in the same regions. Some pits were found in
GaInN QW Sample A1, which are typically observed in
GaInN QWs. Also, even though some undulation in the CL
image of GaInN QW Sample A1 was observed due to the
In inhomogeneity, the CL image showed reasonably highintensity. In contrast, AlGaInN QW Sample A2 showed
many defects on the surface. The inset of the AFM image in
Fig. 1(b) revealed that each defect was surrounded with atrench, and the surrounded region was slightly higher thanthe other ﬂat regions. The height and width of the region

(a)

(b)

Fig. 1. (Color online) AFM, SEM, and CL images of (a) Ga0:90 In0:10 N
QWs (Sample A1) and (b) AlGaInN QWs (Sample A2).

surrounded by the trench are about 4–13 and 60–200 nm,
respectively. In addition, the CL image revealed that theareas of the trench defects were very dark, meaning that the
trench defects should act as nonradiative centers. The trench
defects are not typically observed in GaInN QWs, so thedefects should result from the additional TMAl ﬂow.
Next, to understand the origin of the trench defects,
additional AlGaInN QW samples (B1 to B3) were grown
and characterized. In the AlGaInN samples, various TMIn
ﬂows including zero ﬂow were selected during the AlGaInN
QW growths. Figures 2(a)–2(c) show the surface morphologies of AlGaInN QW Samples B1, B2, and B3 measured by
AFM. As the TMIn ﬂow was reduced, the trench defects
were clearly reduced. Actually, no trench defects wereobserved in AlGaN QW Sample B3, which was grownwithout TMIn ﬂow. Therefore, the origin of the trench
defects in this case should be related to the coexistence of Al
and In at least.
We then grew QW samples at a lower growth pressureof 150 mbar. The purpose of using a lower growth pressure
is to minimize potential parasitic reactions between TMAl,
TMIn, and possibly NH3 by increasing the ﬂow velocity in
the MOVPE reactor ﬂow channel. Figures 3(a) and 3(b)show AFM, SEM, and CL images of GaInN QW Sample C1
and AlGaInN QW Sample C2 grown at 150 mbar, respectively. Clearly, AlGaInN QW Sample C2 shows almost no
trench defects on the surface. Furthermore, a bright CL
image was also obtained from AlGaInN QW Sample C2.
This is attributed to the elimination of the trench defects on
the surface. Another ﬁnding is that the size of the brightregions in the CL image of AlGaInN QW Sample C2[Fig. 3(b)] was smaller than that of those in GaInN QW

08JB27-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JB27

T. Suzuki et al.

Fig. 2. (Color online) AFM images of (a) AlGaInN QWs (Sample B1), (b) AlGaInN QWs (Sample B2), and (c) Al0:07 Ga0:93 N QWs (Sample B3).

Fig. 4. (Color online) PL spectra of AlGaInN QWs (Samples A2 and C2)
and reference GaInN QWs (Samples A1 and C1).

(a)

(b)

Fig. 3. (Color online) AFM, SEM, and CL images of (a) Ga0:87 In0:13 N
QWs (Sample C1) and (b) AlGaInN QWs (Sample C2).

Sample C1 [Fig. 3(a)]. This should be one of the inﬂuencesof adding Al atoms in the well layers, but the details are
unknown at this moment.
PL spectra of AlGaInN QW Samples A2 and C2, andreference GaInN QW Samples A1 and C1 are plotted inFig. 4. AlGaInN QW Sample C2 grown at 150 mbar clearly
showed a drastic improvement of PL intensity comparedwith AlGaInN QW Sample A2, and almost comparable PL
intensities to those of GaInN QW Samples A1 and C1. Also,
the emission peak wavelength of AlGaInN QW Sample C2
showed a blue-shift from that of GaInN QW Sample C1,
resulting in 385 nm emission. This result implies that some
Al atoms incorporated in the well layer contributed to theblue-shift.
Additionally, for an extreme case with the coexistence of
Al and In, AlInN QW Sample C3 was grown at 150 mbar.
As shown in Fig. 5, no trench defects on the surface were

Fig. 5. (Color online) AFM image of Al0:86 In0:14 N QWs (Sample C3).

observed, so the low-pressure growth is conﬁrmed to be a
key technique for obtaining high-quality AlGaInN layers by
suppressing the formation of trench defects. At this moment,
we do not have direct evidence explaining why the lowpressure growth is eﬀective for suppressing the trench defectformation. Further investigations will be necessary.
4. Summary

We have grown AlGaInN QWs based on the growth
conditions of GaInN QWs under diﬀerent growth pressures,
400 and 150 mbar. When the AlGaInN QWs were grown
at 400 mbar, a considerable amount of trench defects were

08JB27-3# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JB27

T. Suzuki et al.

observed on the surface, which acted as nonradiative centers.
The number of trench defects decreased with the decreasing
amount of TMIn ﬂow. When the AlGaInN QWs were grown
at 150 mbar, AlGaInN QWs without trench defects wereobtained, resulting in a high PL intensity with a 385 nm peakwavelength. We found that the low-pressure growth is a
key technique for obtaining high-quality AlGaInN QWs by
suppressing the trench defect formation.

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This study was supported by the MEXT-Supported Program
for the Strategic Research Foundation at Private Universities, 2012–2016.

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GaInN-based tunnel junctions with graded layers

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Applied Physics Express 9, 081005 (2016)
http://doi.org/10.7567/APEX.9.081005

GaInN-based tunnel junctions with graded layers
Daiki Takasuka1, Yasuto Akatsuka1, Masataka Ino1, Norikatsu Koide1, Tetsuya Takeuchi1*,
Motoaki Iwaya1, Satoshi Kamiyama1, and Isamu Akasaki1,2
1Faculty of Science and Engineering, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering, Akasaki Research Center, Nagoya University, Nagoya 464-8601, Japan2*E-mail: take@meijo-u.ac.jp
Received June 20, 2016; accepted July 7, 2016; published online July 22, 2016
We demonstrated low-resistivity GaInN-based tunnel junctions using graded GaInN layers. A systematic investigation of the samples grown bymetalorganic vapor phase epitaxy revealed that a tunnel junction consisting of a 4 nm both-sides graded GaInN layer (Mg: 1 ' 1020 cm%3) and a
2 nm GaN layer (Si: 7 ' 1020 cm%3) showed the lowest speciﬁc series resistance of 2.3 ' 10%4 Ω cm2 at 3 kA/cm2 in our experiment. The InN mole
fraction in the 4 nm both-sides graded GaInN layer was changed from 0 through 0.4 to 0. The obtained resistance is comparable to those ofstandard p-contacts with Ni/Au and MBE-grown tunnel junctions. © 2016 The Japan Society of Applied Physics

he ﬁrst GaInN-based tunnel junction for current
injection from n-layers to p-layers was demonstrated
in 2001.1) Because the tunnel junctions made ofother III–V semiconductors have been adopted in various
devices,2–4) such a GaInN-based tunnel junction is also
expected to be utilized to improve device performance in notjust tandem solar cells5) and cascaded LEDs,6–12) but also
in micro-LEDs,13,14) their arrays,15) lasers,16,17) and verticalcavity surface emitting lasers (VCSELs).6,18) At that time,
however, the voltage drop at the GaInN tunnel junction was
high, about 0.6 V at 50 A=cm2, and further reduction in tunnel
junction resistance was demanded.
In 2010, a new concept to leverage polarization charges at
the GaInN=GaN interfaces was proposed towards realizing
low-resistivity GaInN-based tunnel junctions.19) Thus, theGaInN layer has been used in the tunnel junctions becauseof not only a narrower band gap20) with a lower ionizedenergy of a Mg acceptor 21) but also the polarization doping.
Since then, the GaInN-based tunnel junctions have shownmarked improvements in their reverse-biased characteristics.
In particular, a GaInN tunnel junction grown by MBE
showed a very low speciﬁc series resistance of 1.2 × 10−4
Ω cm2.22) On the other hand, a GaInN tunnel junction grownby metalorganic vapor phase epitaxy (MOVPE) showed aresistance of 4.0 × 10−4 Ω cm2 by using a very high InN mole
fraction of 0.4 and lateral Mg activation.23–25) However, alarge voltage drop at the GaInN tunnel junction of about
1 V was still observed under operation with a high current
density of ∼3 kA=cm2.
In this study, we paid attention to not only a tunnelingprobability component in the GaInN layer but also an energy
spike component at the GaInN=GaN interfaces in terms of theresistance of the GaInN tunnel junction. The GaInN layer
with the high InN mole fraction leads to a high tunnelingprobability owing to the small band gap, but at the same time
generates large energy spikes at the GaInN=GaN interfaces,
which result in a high resistance. Such energy spikes arenotorious as a cause of the large resistance in semiconductorbased devices, such as distributed Bragg reﬂectors. In order
to eliminate such energy spikes, graded layers were proposed
and often used.26,27) Actually, the eﬀect of the graded AlGaN
layers in a GaN=AlN tunnel junction was theoretically investigated by Schubert.28) Here, we experimentally investigated
the impact of the graded GaInN layers on the resistance of the
GaInN tunnel junctions. Various InN mole fraction proﬁles inT

Fig. 1. Sample structure consisting of a standard blue LED structure,
a GaInN-based tunnel junction, a 400-nm top n-GaN layer, and a 10-nm
n+-GaN contact layer.

the graded GaInN tunnel junctions were also determined.
As a result, we demonstrated a graded GaInN tunnel junction
with a very low resistivity, the speciﬁc series resistance of
which was 2.3 × 10−4 Ω cm2 at 3 kA=cm2.
All the samples were grown by MOVPE on c-planesapphire substrates using low-temperature-deposited buﬀer
layers.29) As shown in Fig. 1, a sample structure consistedof a standard blue LED structure, a GaInN-based tunnel
junction, a 400 nm top n-GaN layer, and a 10 nm n+-GaN
contact layer. We prepared four samples containing diﬀerent
GaInN layers in the tunnel junctions. Figure 2 shows thedesigned InN mole fraction proﬁles in the GaInN layers. The
ﬁrst sample shown in Fig. 2(a) contained a 2 nm2 Ga0.6In0.4N
layer, which was the same layer that we previously optimized.24) The other three samples contained graded GaInN
layers. The second sample shown in Fig. 2(b) had a 4 nm
p-side graded GaInN layer in which the InN mole fraction
was increased from 0 at the p-side (left) to 0.4 at the n-side(right). The third sample included the opposite n-side graded
GaInN layer in which the InN mole fraction was decreased
from 0.4 at the p-side to 0 at the n-side, as shown in Fig. 2(c).
Finally, as shown in Fig. 2(d), the fourth sample contained a
4 nm both-sides graded GaInN layer in which the InN mole
fraction was changed from 0 to 0.4 in a 2 nm p-side GaInN
layer and from 0.4 to 0 in a 2 nm n-side GaInN layer. In other081005-1© 2016 The Japan Society of Applied Physics

Appl. Phys. Express 9, 081005 (2016)

D. Takasuka et al.

(a)

(b)

(c)

(d)

Fig. 2. Designed InN mole fraction proﬁles in the GaInN tunnel junctions: (a) a 2 nm2 Ga0.6In0.4N, (b) a 4 nm p-side graded GaInN, (c) a 4 nm n-side graded
GaInN, and (d) a 4 nm both-sides graded GaInN.

(a)

(b)

(c)

(d)

Fig. 3. Calculated band diagrams around the GaInN tunnel junctions: (a) a 2 nm2 Ga0.6In0.4N, (b) a 4 nm p-side graded GaInN, (c) a 4 nm n-side graded
GaInN, and (d) a 4 nm both-sides graded GaInN.

words, the InN mole fraction proﬁle of the both-sides graded
GaInN layer was “Λ”-shaped. Here, the InN mole fractions
were controlled by changing the TMIn ﬂow rate during thegraded GaInN layer growth while the growth temperature
and TMGa ﬂow rate remained constant. Note that the total
amount of In in the tunnel junction was designed to be thesame among all the samples. All the GaInN layers in thetunnel junctions were doped with Mg of 1 × 1020 cm−3. Then,
2 nm heavily Si-doped (7 × 1020 cm−3) GaN layers weregrown on the GaInN layers to complete the tunnel junction
structures.
A lateral Mg activation previously developed for theLEDs with the tunnel junctions was performed for all thesamples.25) After forming device mesas of 35 µm diameter,
the wafers were thermally annealed under oxygen ambientat 725 °C for 30 min. Metal layers, Ti (30 nm)=Al (100 nm)=Ti (20 nm)=Au (150 nm), for the top and bottom n-contacts
were simultaneously formed. Since the small mesas wereformed, we were able to measure the current density–voltage( j–V ) and current density–light output power ( j–L) characteristics of the LEDs with the tunnel junctions under a highcurrent density of up to 5 kA=cm2. For comparison, we also
prepared a small standard blue LED with a p-contact with
Ni=Au.
Figure 3 shows calculated band diagrams around thetunnel junctions in the four samples. Commercial software(SiLENSe)30) was used for the calculation. Large energy
spikes were clearly shown at the interfaces of the 2 nm2GaInN layer in Fig. 3(a). The energy heights of the spikes
were 0.9 eV at the conduction band and 0.4 eV at the valenceband. On the other hand, no energy spikes were shown at thep-side interface in the p-side graded GaInN tunnel junction

081005-2© 2016 The Japan Society of Applied Physics

Appl. Phys. Express 9, 081005 (2016)

D. Takasuka et al.

(a)Fig. 4. j–V characteristics of the four LEDs with the GaInN tunnel
junctions containing the 2 nm2 Ga0.6In0.4N, 4 nm p-side graded GaInN, 4 nm
n-side graded GaInN, or 4 nm both-sides graded GaInN.

(b)[Fig. 3(b)] and at the n-side interface in the n-side graded
GaInN tunnel junction [Fig. 3(c)] owing to the graded layers.
Interestingly, even with the very thin graded layers (2 nm)[Fig. 3(d)], a clear reduction in the number of energy spikes
was also shown at both interfaces in the both-sides graded
GaInN tunnel junction.28)
Figure 4 shows the j–V characteristics of the four LEDs
with the square GaInN tunnel junctions or various graded
GaInN tunnel junctions. The operating voltage of the LEDwith the square GaInN tunnel junction was 5.3 V at 3 kA=cm2.
On the other hand, all of the LEDs with the graded GaInN
tunnel junctions showed lower operating voltages of 4.6 V orlower at 3 kA=cm2. The diﬀerential speciﬁc series resistanceof the device was also estimated to be 2.3 × 10−4 Ω cm2 at
3 kA=cm2, suggesting that the tunnel junction resistanceshould be lower than this value, which is comparable to thebest value of MBE-grown nitride-based tunnel junctions.22)
These lines of evidence suggest that the energy spikes at theheterointerfaces have a considerable impact on the tunnel
junction resistances. Note that the LED with the both-sides
graded tunnel junction showed the lowest voltage of 4.3 V
at 3 kA=cm2, but the value was not so low as expected inconsideration of the results obtained from the single-side(the p- or n-side) graded tunnel junctions. One of the possible
reasons is that the actual 2 nm graded GaInN layer could be
too thin to eliminate the energy spike suﬃciently.
Then, we attempted a larger graded layer thickness inthe both-sides graded GaInN tunnel junction. We found thatthe sample surface deteriorated when we grew an 8 nm bothsides graded GaInN layer from 0 through 0.4 to 0 on the LED
structure. It is diﬃcult to obtain simultaneously a small band
gap and thick graded layers with high InN mole fractions
owing to the large lattice mismatch between GaInN and GaN.
In other words, the choice of InN mole fraction in the graded
GaInN tunnel junction aﬀects two opposite aspects. One is
the tunneling probability in which a “higher” InN mole fraction results in a lower resistance, and the other is the energy
spike height in which a “lower” InN mole fraction results in alower resistance. Therefore, an optimum InN mole fraction
may exist in such a trade-oﬀ situation. We then prepared
two additional LEDs with diﬀerent both-sides graded GaInN
tunnel junctions. As shown in Fig. 5, two diﬀerent InN mole
fraction proﬁles were obtained. The ﬁrst one was a 6 nm
both-sides graded GaInN layer with InN mole fractions from0 through 0.35 to 0, as shown in Fig. 5(a). The second one

Fig. 5. Designed InN mole fraction proﬁles in the graded GaInN tunnel
junctions: (a) a 6 nm both-sides graded Ga0.65In0.35N and (b) an 8 nm
both-sides graded Ga0.8In0.2N.

(a)

(b)
Fig. 6. Calculated band diagrams around the graded GaInN tunnel
junctions: (a) a 6 nm both-sides graded Ga0.65In0.35N and (b) an 8 nm
both-sides graded Ga0.8In0.2N.

was an 8 nm both-sides graded GaInN layer with InN mole
fractions from 0 through 0.2 to 0, as shown in Fig. 5(b).
Figure 6 shows calculated band diagrams around thetunnel junctions in the two samples. No energy spikes were
found at both interfaces in the two samples. Actually, as
compared with the above-mentioned 4 nm both-sides graded
GaInN tunnel junction in Fig. 3(d), the additional two
samples containing the wider graded layer with the lower
InN mole fraction showed smoother potential proﬁles but
slightly larger depletion lengths. Figure 7 shows the j–V
characteristics of the two additional both-sides graded GaInN
tunnel junctions (6 and 8 nm) as well as the 4 nm both-sides
graded GaInN tunnel junction. In this experiment, we foundthat the 4 nm graded GaInN layer with the 0.4 InN mole
fraction still resulted in the lowest resistance.

081005-3© 2016 The Japan Society of Applied Physics

Appl. Phys. Express 9, 081005 (2016)

D. Takasuka et al.
Exploratory Research (No. 15K13959), and the Ministry of Education, Culture,
Sports, Science and Technology (MEXT)-Supported Program for the Strategic
Research Foundation at Private Universities, 2012–2016.

Fig. 7. j–V characteristics of the three LEDs with the tunnel junctions
containing the 4 nm both-sides graded Ga0.6In0.4N, 6 nm both-sides graded
Ga0.65In0.35N, or 8 nm both-sides graded Ga0.8In0.2N.

Fig. 8. j–V–L characteristics of the LED with the 4 nm both-sides graded
Ga0.6In0.4N tunnel junction and a standard LED with p-contact and Ni=Au.

Eventually, we plotted the j–V–L characteristics of the
LED with the 4 nm both-sides graded Ga0.6In0.4N tunnel
junction showing the lowest resistance and the standard LED
containing the p-contact with Ni=Au. Here, light output
power was measured under the bottom emission conﬁguration. As shown in Fig. 8, the operating voltage and light
output power of the LED with the tunnel junction werecomparable to those of the standard LED with p-contact.
In this study, we investigated the impact of the graded
GaInN layers on the resistance of the GaInN tunnel junctions
grown by MOVPE. The experimental results revealed thatthe graded GaInN layers resulted in lower voltage drops at
the tunnel junctions. Our lowest speciﬁc series resistance ofthe graded GaInN tunnel junction was 2.3 × 10−4 Ω cm2 at
3 kA=cm2 as obtained by using the 4 nm both-sides graded
Ga0.6In0.4N layer. The value is now comparable to those of
the standard p-contact with Ni=Au and the GaInN tunnel
junction grown by MBE. Such a low-resistivity GaInN tunnel
junction grown by MOVPE could open up great possibilities
for high-performance nitride-based optoelectronic devices.
Acknowledgments This work was supported by a Grant-in-Aid for
Scientiﬁc Research (B) (No. 26286045), a Grant-in-Aid for Challenging

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30) Web [http://www.str-soft.com/products/SiLENSe/index.htm].

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Electrical properties of n-type AlGaN with high Si concentration

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REGULAR PAPER

Japanese Journal of Applied Physics 55, 05FE02 (2016)
http://doi.org/10.7567/JJAP.55.05FE02

Electrical properties of n-type AlGaN with high Si concentration
Kunihiro Takeda1*, Motoaki Iwaya1, Tetsuya Takeuchi1, Satoshi Kamiyama1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan2*E-mail: 143434025@ccalumni.meijo-u.ac.jp
Received November 5, 2015; accepted December 1, 2015; published online March 30, 2016
The electrical properties of Si-doped AlGaN layers (AlN molar fractions: 0.03–0.06) with the donor concentrations (ND) from 8.8 ' 1017 to
4.5 ' 1020 cm%3 were investigated by variable-temperature Hall effect measurement using the van der Pauw method. A minimum resistivity of
3.6 ' 10%4 Ω cm was obtained for Si-doped AlGaN with a smooth surface at room temperature. We found that the activation energy of the Si donor
is affected by the Coulomb interaction in the AlGaN layer with ND values from 8.8 ' 1017 to 2.5 ' 1020 cm%3. In several AlGaN layers, the freeelectron concentration did not vary with sample temperature, as expected in the case of degeneracy. The localization of GaN in the AlGaN layer
was speculated as a cause of degeneracy of samples. © 2016 The Japan Society of Applied Physics

1.

Introduction

High-performance AlGaInN-based green, blue, and UV lightemitting diodes (LEDs) are commercialized worldwide.1–5)
Nowadays, the external quantum eﬃciencies (EQEs) of blueand near-UV LEDs are already more than 80%4) and 40%,5,6)
respectively. From a practical viewpoint, the improvement ofwall plug eﬃciency (WPE) in these LEDs is essential. A
decrease in device resistance including the sheet and contact
resistances of n- and p-type layers is important for improving
WPE.7) A reduction in resistivity (ρ) in an n-type layer isimportant, because the layer has a role in current spreading
and contact layers. Many studies have reported on the
reduction of ρ in n-type GaN.8–11) Fritze et al. realized ann-type GaN layer with a high free electron concentration (ne)(∼2.9 × 1020 cm−3) and low ρ (∼6.3 × 10−4 Ω cm) at RT by
Ge doping.10) Since Si is often used as an n-type dopant in
nitride-based LED device manufacturing, low-ρ Si-dopedn-type layers are preferred. We have reported that high ne(∼1.4 × 1020 cm−3) and low-ρ (∼5.9 × 10−4 Ω cm) Si-doped
Al0.05Ga0.95N ﬁlms have been obtained at RT with an
atomically ﬂat surface.7) In addition, violet LEDs fabricatedusing a low-ρ Si-doped n-type Al0.05Ga0.95N layer wererealized to reduce device resistance and improve WPE.7) On
the other hand, a number of studies have shown the activationenergy (ED) of the Si donor in GaN.11–13) In these studies, theelectrical conduction properties of Si-doped GaN with a Siconcentration of approximately 1019 cm−3 have been shownto degenerate, which is deﬁned as the Fermi level exit in the
conduction band. Nevertheless, the detailed electrical properties and minimum ρ of Si-doped AlGaN with a low AlN
molar fraction have not yet been elucidated.
In this study, we obtained the reduction bounds ofresistivity and ED in Si-doped AlGaN layers (AlN molar
fractions: 0.03–0.06) with donor concentrations (ND) from
8.8 × 1017 to 4.5 × 1020 cm−3 and a ﬂat surface by variabletemperature Hall eﬀect measurement. Moreover, the degeneracy of AlGaN layers with high ND was investigated.
2.

Experimental methods

All samples were grown using metal organic vapor-phase
epitaxy (MOVPE) reactors. Figure 1 shows a schematic viewof the sample structure. Trimethylaluminum, trimethylgallium, and ammonia (NH3) were used as the precursors of

Fig. 1. Schematic view of the sample structure.

Al, Ga, and N, respectively. Silane (SiH4) or tetramethylsilane [Si(CH3)4] was used as a dopant source. A low-temperature (LT) buﬀer layer was used as a cover to obtain ahigh-crystalline-quality GaN underlying layer on a c-planesapphire substrate. Thereafter, an unintentionally doped
approximately 3-µm-thick GaN underlying layer and a Sidoped AlGaN (AlN molar fractions: 0.03–0.06) layer weresequentially grown. The thicknesses of AlGaN layers werefrom 1.2 to 1.5 µm. The AlGaN layers were grown at aconstant pressure of 150 Torr. The growth temperature of
AlGaN layers was set to be in the range from 1,050 to
1,080 °C and the V=III ratio was 3,400. The ne, mobility (μe),
and ρ in each sample were characterized by variabletemperature Hall eﬀect measurement using the van der Pauw
method.
3.

Results and discussion

Figure 2 shows the Si=III ratio dependence of the ne and ρ in
Si-doped AlGaN at RT. The Si-doped AlGaN layers weregrown at various Si=III ratios from 6.5 × 10−6 to 5.2 × 10−3.
The maximum ne and minimum ρ of 4.5 × 1020 cm−3 and3.6 × 10−4 Ω cm, respectively, were obtained at a Si=III ratioof 3.9 × 10−3 and at RT. At Si=III ratios from 6.5 × 10−6 to
3.9 × 10−3, the samples were observed to have a smoothsurface by diﬀerential interference contrast microscopy
measurement and also a root mean square (RMS) roughnessof less than 0.31 nm at 5 × 5 µm2 measured by atomic force
microscopy (AFM). The full width at half maximum(FWHM) values of AlGaN(0002) and the ð1012Þ
X-ray

05FE02-1© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FE02 (2016)

K. Takeda et al.

Fig. 2. (Color online) Si=III ratio dependence of ne and ρ at RT.

rocking curve (XRC) ω-scan were very much the same as
those of the GaN underlying layer. FWHMs of approximately
290 arcsec for the AlGaN(0002) ω-scan and 330 arcsec for
the AlGaNð1012Þ
XRC ω-scan were observed in thesesamples. In calculating ED, the donor levels of defects andunexpected impurities were ignored, because ne was almost
proportional to the Si=III ratio in a range from 6.5 × 10−6 to
3.9 × 10−3, as shown in Fig. 2.
In contrast, the decrease in ne and increase in ρ wereobserved in Si-doped AlGaN with a Si=III ratio of
5.2 × 10−3, as shown in Fig. 2. The surface of this AlGaN
layer was quite rough and the FWHM value of AlGaNð1012Þ
XRC ω-scan was increased. The maximum ne in Si-doped
GaN with a smooth surface has been reported to beapproximately 2 × 1019 cm−3 at RT.10) In contrast, we
obtained an n-type layer with a smooth surface and ne of
4.5 × 1020 cm−3 at RT using GaN with addition of asmall amount of AlN. Thus, we investigated the electricalproperties of samples with the ne of less than 4.5 × 1020 cm−3
at RT.
Figures 3 and 4 show the results of variable-temperatureHall eﬀect measurement of each sample. These ﬁgures show
ne as a function of reciprocal temperature. From the 12 grownsamples, 4 showed electrical properties such as degeneracy.
Accordingly, we classiﬁed the samples into two types fromthe temperature dependence of ne. Group I showed ne thatvaries with sample temperature, as shown in Fig. 3. In
contrast, Group II showed ne that did not vary with sample
temperature in the range from 300 to 70 K, as shown in
Fig. 4. This electrical property suggests degeneracy.11–13)
First, we investigated the ED of the Si donor in the AlGaN
layer from the samples in Group I. In Fig. 3, the samples
with ne values of more than 6.0 × 1018 cm−3 at RT showeda variation of ne at approximately RT and ne saturation at
150 K or below. Since the C, O, H, Mg, and Zn impurity
concentrations in the AlGaN layer of each sample were
determined to be below the secondary ion mass spectroscopy
detection limits, the eﬀect of donor impurities excluding
Si was ignored. A two-band model that includes electrontransport within a donor impurity band as well as in the
conduction band should be considered to explain theoccurrence of the variation and saturation of ne.14–16) Thus,
the ED values of the samples with ne values of more than
6.0 × 1018 cm−3 at RT were obtained by ﬁtting the temperature dependence of ne at approximately RT.

Fig. 3. (Color online) ne in Si-doped AlGaN (AlN molar fractions:
0.03–0.06) as a function of reciprocal temperature characterized by thevariable-temperature Hall eﬀect measurement. This ﬁgure only plotted theresults of Group I. Circles and triangles in the ﬁgure indicate the samples
grown in MOVPE reactors using SiH4 and Si(CH3)4, respectively.

Fig. 4. (Color online) ne in Si-doped AlGaN (AlN molar fraction: 0.03) asa function of reciprocal temperature determined by variable-temperature Halleﬀect measurement. This ﬁgure only plotted the results of Group II. The
samples in this ﬁgure were grown in MOVPE reactors using Si(CH3)4.

From Fig. 2, ED was calculated as a single-donor level ofthe Si dopant. When the donor level is regarded as a singleimpurity level, the activation energy of the dopant wascalculated from the charge neutrality condition taking
residual acceptors into consideration. In the charge neutrality
equation for the n-type semiconductor, ED, ND, and NA are
related through the equation11,17–20)
ne ðne þ NA Þ
NC
ED
exp 
¼
;
ð1Þ
ND  NA  ne2kT2me kT 3=2
NC ¼ 2;ð2Þ
h2
where NC is the eﬀective density of states in the conduction
band, k is the Boltzmann constant, T is the temperature, me is

05FE02-2© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FE02 (2016)

K. Takeda et al.

Fig. 5. (Color online) ED of Si-doped AlGaN (AlN molar fractions:
0.03–0.06) as a function of N1=3
D . The ED values were obtained in the range
between 12.0 and 17.0 meV. The ﬁtting line is shown in the ﬁgure. From thisﬁtting line, the inherent donor energy observed at very low donor
concentrations was estimated to be approximately 17.6 meV.

Fig. 6. (Color online) PL spectra from (a) AlGaN layer grown using SiH4and (b) AlGaN layer grown using Si(CH3)4 obtained using a 266 nm third
harmonics YAG laser at RT. The solid and dashed lines indicate samples ofGroup I and Group II, respectively.

the electron eﬀective mass, and h is the Planck constant. Theme in AlGaN of each sample were presumed by20,21)

grown using Si(CH3)4 as the Si source; therefore, thediﬀerence in the Si source is considered as the cause of thediﬀerent electrical properties occurring in AlGaN layers of
the same ne.
Figure 6 shows the photoluminescence (PL) spectra
obtained using a 266 nm third harmonics yttrium aluminum
garnet (YAG) laser at RT. Figures 6(a) and 6(b) show thePL spectra of the AlGaN layer grown using SiH4 andSi(CH3)4, respectively. In Fig. 6, the solid line and dashedline indicate PL spectra of nondegenerate and degeneratesamples, respectively.
As shown in Fig. 6(b), the spectral patterns of the samples
grown using Si(CH3)4 showed that the peaks of AlGaN andapproximately 365 nm in degenerate samples with ND valuesof 2.8 × 1019 and 4.5 × 1020 cm−3, despite having approximately the same thickness. A particularly large change wasobserved in the sample with an ND of 4.5 × 1020 cm−3. On theother hand, as shown in Fig. 6(a), even when ND increased,
the spectral patterns of the samples grown using SiH4 did notshow a peak of approximately 365 nm. From these diﬀerent
spectral patterns, the localization of GaN was observed indegenerate samples.
Figure 7 shows the Si-doped GaN layer with ne of
9.0 × 1018 cm−3 as a function of reciprocal temperature
determined by the variable-temperature Hall eﬀect measurement. This GaN layer was grown in the MOVPE reactor
using SiH4. Similarly to Group II, this Si-doped GaN layer
showed degeneracy. Regarding the results, Götz et al. havealready reported that the ne of Si-doped GaN with a Siconcentration of 2 × 1019 cm−3 does not vary with sample
temperature, indicating a doping level above the degeneracy
limit.11) From our results and the report,11) the electricalproperties of Si-doped GaN with ND of more than 9.0 ×
1018 cm−3 should change degeneracy.27)
The localization of GaN in the AlGaN layer can beconsidered as the cause of degeneracy. Because the bandgap
energy of AlGaN with an AlN molar fraction of 0.03 ishigher by more than 80 meV than that of GaN, the AlGaN
layer with the ND of approximately 2 × 1019 cm−3 would not
attain the degeneracy. The electrical properties of AlGaNme ¼ ðm2e?  mek Þ1=3  m0 :ð3Þ

mek and me? were 0.20 and 0.18 in GaN, and 0.33 and 0.25
in AlN, respectively.18,21) We obtained ED by ﬁtting using
Eq. (1). Fit parameters were ED, ND, and NA.
Götz et al. have suggested a linear function for theacceptor activation energies (EA) in Mg-doped GaN.22) From
this case, the donor activation energy (ED) and EA should be
described as17–19,23,24)
ED,AðND,A Þ ¼ ED,Að0Þ  f

q2 1=3N ;
4" D,A

ð4Þ

where f is a geometric factor of value Γ(2=3)(4π=3)1=3, q is
the electronic charge, ε is the dielectric constant, ED(0) is an
inherent donor energy that corresponds to hydrogen-like
donors observed at very low ND,24) and NA is the acceptor
concentration. From Eq. (4), the ED,A of the impurity
semiconductor is aﬀected by Coulomb interaction between
band carriers and ionized impurities. From this theory, ED,A ispresumed to be proportional to the 1=3 power of ND,A.
Several impurity semiconductors have been fabricated on thebasis of this theory;22,24–26) nevertheless, there has been noreport in n-type nitride semiconductors.
Figure 5 shows a summary of the ED of the Si donor in
AlGaN layers as a function of ND. ED decreased fromapproximately 17.0 to 12.0 meV with increasing ND from
8.8 × 1017 to 2.5 × 1020 cm−3. ED(0) was estimated to beapproximately 17.6 meV by ﬁtting in Fig. 5. From a ﬁtting
line, the ED of Si-doped AlGaN with a low AlN molar
fraction was conﬁrmed to vary proportionally to the 1=3
power of ND. Therefore, we concluded that the ED of Sidoped AlGaN (AlN molar fractions: 0.03–0.06) is aﬀected by
Coulomb interaction between conduction band electrons andionized donors.
Next, we investigated the cause of electrical properties
such as degeneracy in Group II. The samples with ne valuesof more than 3.3 × 1019 cm−3 at RT in Group I were grownusing SiH4 as the Si source. All the samples in Group II were

05FE02-3© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FE02 (2016)

K. Takeda et al.

8.8 × 1017 to 2.5 × 1020 cm−3, ED(0) was also estimated to beapproximately 17.6 meV. We found that the ED of the Si
donor is aﬀected by the Coulomb interaction in the AlGaN
layer with ND values from 8.8 × 1017 to 2.5 × 1020 cm−3. In
several AlGaN layers with high ND values, ne did not varywith sample temperature, as expected in the case ofdegeneracy. The cause of the degeneracy of samples was
assumed to be the localization of GaN in the AlGaN layer.
Acknowledgments

Fig. 7. (Color online) ne in Si-doped GaN layer with ne of 9.0 ×
1018 cm−3 as a function of reciprocal temperature determined by variabletemperature Hall eﬀect measurement. Similarly to Fig. 3, this Si-doped GaN
layer showed degeneracy.

and GaN can be obtained when GaN forms partially in the
AlGaN layer. When GaN localizes in AlGaN, the sample
with high ND might show degeneracy in GaN, even whennondegenerate in AlGaN. Since a peak of approximately
365 nm was observed in degenerate samples, as shown in
Fig. 6(b), this assumption is possible. From the diﬀerence inluminescence intensity at approximately 365 nm in the PL
spectra of the ND values of 2.8 × 1019 and 4.5 × 1020 cm−3,
there is a possibility that the localization of GaN is enhancedwith the increase in Si(CH3)4 supply. The samples weregrown in two types of MOVPE reactor using Si(CH3)4.
AlGaN layers with ND values of more than 2.8 × 1019 cm−3
grown in each MOVPE reactor showed degeneracy. In
particular, the Si-doped Al0.09Ga0.91N layer with ND of 4.8 ×
1019 cm−3 in unintentionally doped AlN underlying layergrown in the MOVPE reactor using Si(CH3)4 was conﬁrmed
to shown degeneracy, and the PL spectral pattern measured
using the 266 nm third harmonics YAG laser at RT showedthat the peaks of AlGaN and approximately 365 nm.
The localization of GaN might tend to occur when using
Si(CH3)4 as a dopant source. However, we are unable toconclude that the localization of GaN in AlGaN is the causeof the degeneracy of samples with AlGaN layers with an AlN
molar fraction of 0.03 and ND values of more than 2.8 ×
1019 cm−3. Therefore, further studies are needed to determine
the cause of degeneracy of samples.
4.

Conclusions

We grew Si-doped AlGaN layers with ND values from
8.8 × 1017 to 4.5 × 1020 cm−3 using MOVPE reactors with
SiH4 or Si(CH3)4 as a dopant source. The minimum-ρ (3.6 ×
10−4 Ω cm) was obtained at the Si-doped AlGaN layer with
the ne of 4.5 × 1020 cm−3 and a ﬂat surface. The samples wereclassiﬁed into two types from the temperature dependence ofne by the variable-temperature Hall eﬀect measurement. Fromthe sample group in which ne varies with sample temperature,
the ED of the Si donor in AlGaN layers decreased fromapproximately 17.0 to 12.0 meV with increasing ND from

This study was partially supported by the Program for
the Strategic Research Foundation at Private Universities,
2012–2016, supported by the Ministry of Education, Culture,
Sports, Science and Technology (MEXT), a MEXT Grantin-Aid for Specially Promoted Research No. 25000011, anda MEXT Grant-in-Aid for Scientiﬁc Research A No.
15H02019.

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Metal–Organic Vapor Phase Epitaxy Growth of Embedded Gallium Nitride Nanocolumn for
Reduction in Dislocation Density

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2013 Jpn. J. Appl. Phys. 52 08JE23(http://iopscience.iop.org/1347-4065/52/8S/08JE23)
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Theeradetch Detchprohm, Masahiro Yano, Shigekazu Sano et al.

REGULAR PAPER

Japanese Journal of Applied Physics 52 (2013) 08JE23
http://dx.doi.org/10.7567/JJAP.52.08JE23

Metal–Organic Vapor Phase Epitaxy Growth of Embedded Gallium Nitride Nanocolumn
for Reduction in Dislocation Density
Shinya Umeda1 , Takahiro Kato1 , Tsukasa Kitano2 , Toshiyuki Kondo2 , Hiroyuki Matsubara1 ,Satoshi Kamiyama1 , Tetsuya Takeuchi1 , Motoaki Iwaya1 , and Isamu Akasaki1;31Department of Materials Science and Engineering, Meijo University, Nagoya 468-0073, Japan
El-seed, Nagoya 468-0073, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
E-mail: 123434006@ccalumni.meijo-u.ac.jp
2

Received October 14, 2012; accepted May 19, 2013; published online August 20, 2013
The use of nanocolumn crystals is thought to be effective in producing a low-dislocation-density GaN layers. In this paper, we propose a metal–organic vapor phase epitaxial (MOVPE) growth method for producing uniform GaN nanocolumns using deep through-holes in a thick SiO2
selective growth mask. A SiO2 film with a thickness of 500 nm was deposited by sputtering on an AlN buffer layer/SiC substrate. A nanoimprinting
technique was applied to produce dot openings. Then, dry etching with CF4 gas was carried out to form deep through-holes in the SiO2 film. In thesecond MOVPE growth, individual GaN nanocolumns coalesced into a planarized GaN layer, after thinning the SiO2 mask to 100 nm. A cathodeluminescence image of the GaN layer on a GaN nanocolumn template shows a low dislocation density of 1:3  108 cm 2 , while that of a GaN layer
directly grown on an AlN buffer layer shows a dislocation density of 9:4  108 cm 2 . # 2013 The Japan Society of Applied Physics

1. Introduction

Nitride semiconductor nanocolumns are one-dimensional
nanoscale crystals without dislocations. Such crystals arecolumnar materials with diameters of about several tens
to several hundreds of nm, where threading dislocations
penetrating along the growth direction bend horizontally andare terminated in the process.1) Owing to this advantage,studies of high-performance light-emitting devices with
nanocolumn structures have been focused on.2,3) Nanocolumn materials are also promising for the growth of a highquality nitride layer by nanoscale epitaxial lateral overgrowth (nano-ELO).4) Nano-ELO can signiﬁcantly reduce
the epilayer thickness for planarization, compared withconventional microscale ELO. Therefore, this approach may
give a great advantage for GaN growth on foreign substrates
such as Si and SiC,5–7) which have a thermal expansion
mismatch with GaN and a thickness limitation due to crack
formation. Moreover, the selective growth of alloys, suchas AlGaN and InGaN, which is impossible in conventional
microscale ELO, may be possible, because nano-ELO
requires only a short migration distance of species on thegrowth surface. Similarly to the conventional microscale
ELO, the use of nanocolumns may also allow heteroepitaxial
growth without the formation of misﬁt dislocations andcracks, because the stress caused by lattice and thermal
expansion mismatches at a heterointerface can be relaxed.
Nanocolumn crystals are prepared mainly by molecularbeam epitaxial (MBE) growth,8,9) and only a limited numberof reports on the metal–organic vapor phase epitaxial(MOVPE) growth of nanocolumns have been published.10–14) This may be due to the fact that the MOVPE
growth of nanocolumns makes it diﬃcult to control the inplane uniformity under small-scale selective growth. Actually, there are no reports on the selective-area growth ofnanocolumns by MOVPE, in terms of uniform growth overa wide area of several cm2 . In our experiments using thisapproach, the uniformity of nanocolumns over a 2-in. wafer
has not yet been achieved. Considering their device
application, it is indispensable to grow uniform nanocolumns in a wide area.

In this study, we demonstrate a method of high-quality
GaN layer growth on SiC, which is composed of deep
through-holes in a SiO2 ﬁlm, embedded GaN nanocolumns,
and a planarized GaN overlayer. The method of embedding
nanocolumns deep into SiO2 ﬁlms makes it possible to showa high structural uniformity in a wafer with a wide area ofover 2 in. In addition, the GaN layer grown on such nanocolumns with a low dislocation density of 1:3  108 cm 2
has been achieved.
2. Experimental Procedure

Figure 1 shows the procedure for GaN nanocolumn growthby MOVPE. First, a 100-nm-thick AlN buﬀer layer was
grown on a 2-in. SiC substrate. The AlN buﬀer layer iscommonly used for growing a high-quality GaN crystal
on a SiC substrate.15–17) A SiO2 ﬁlm with a thickness of 30
or 550 nm was then deposited by RF sputtering on thissubstrate. Nanoimprinting lithography was applied to thepatterning of nanocolumns. A thermally cured resist was
spin-coated on the substrate, and the pattern was transferredfrom a mold having pillars in a triangular arrangementwith a pitch of 460 nm and a diameter of 250 nm to thesubstrate. Then, dry etching by CF4 gas was carried out toform through-holes in the SiO2 ﬁlm. After removing thenanoimprinting resist, GaN nanocolumns were grown at atemperature of 950  C, a V/III ratio of 1000, and a pressureof 100 mbar, which seem to be suitable for nanocolumn

growth. These growth conditions tend to form {1101}
facets.
The two substrates with patterned 30 and 550-nm-thick
SiO2 ﬁlms were simultaneously set in the MOVPE reactor,
and GaN nanocolumns were grown under the same growth
conditions. In the case of the 30-nm-thick SiO2 ﬁlm, atypical selective growth of GaN nanocolumns can beachieved.
The SiO2 mask was then etched down to a thickness of
100 nm to expose the upper part of the nanocolumns in orderto promote the lateral growth and planarization of the GaN
layer. When the etching of SiO2 was not achieved, large
island crystals were grown, and as a result, the planarizationof the overgrown GaN layer became diﬃcult. On the other
hand, a planarized GaN layer was easily obtained in the

08JE23-1# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JE23

S. Umeda et al.

Regrown GaN

GaN

resist
SiO2
AlN
SiC

Fig. 1. (Color online) Growth procedure of GaN nanocolumns and regrowth of planarized GaN layer.

SiO2
AlN
SiC
1.0µm

1.0µm

Fig. 2. SEM images of top and cross-sectional views of SiO2 mask with deep through-holes.

(a) top view(b) cross-sectional view(c) bird’s-eye viewThin film
mask
1.0µm

1.0µm

1.0µm

1.0µm

1.0µm

1.0µm

Thick film
maskFig. 3. SEM images of (a) top, (b) cross-sectional, and (c) bird’s-eye views of GaN nanocolumns grown on thin and thick SiO2 masks. For the observationof the bird’s-eye view SEM image, the SiO2 was removed.

case that all the SiO2 was removed before the regrowth of
GaN. However, no reduction in dislocation density wasobserved. Therefore, such a process using of SiO2 masks
suitable thickness is required for the planarized growth of
GaN. After that, the regrowth of the GaN layer was
performed on a GaN nanocolumn template. For comparison,
the regrowth of the GaN layer was also carried out on a ﬂat
AlN/SiC substrate. The growth conditions of a temperatureof 1050  C, a pressure of 100 mbar, and a V/III ratio of2700 were used for the planarized growth because of theenhancement of lateral growth.18) Dislocation densities were
measured using the cathodoluminescence (CL) mapping
images of these samples.
3. Results and Discussion

Figure 2 shows scanning electron microscopy (SEM) imagesof the top and cross-sectional views of the SiO2 mask after
achieving the through-hole arrangement. Uniform through-holes were successfully obtained by our process, although
their sidewalls were not perfectly smooth. From the images,
a period of 460 nm and a hole diameter of 200 nm wereconﬁrmed. The etching of the SiO2 mask was accurately
terminated on the surface of the AlN buﬀer layer, and thedepth of the through-holes was 550 nm. Figure 3 shows
SEM images of the top, cross-sectional, and bird’s-eye viewsof GaN nanocolumns. In the thin (30 nm) SiO2 mask, small
pyramids and large islands were randomly observed. Owing
to the large diﬀusion length of Ga-containing species on thesurface, large islands tended to form, interrupting nanocolumn growth. Under typical MOVPE growth conditions, thesupply rate of precursors was considered to be very highagainst a small volume of nanocolumn crystals. Special
techniques, such as the pulsed supply of precursors,13,14)
may be indispensable in the MOVPE growth of nanocolumns. On the other hand, uniform nanocolumns with aheight of 600 nm were obtained in through-holes in the thick

08JE23-2# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JE23(a)

S. Umeda et al.

(b)regrown GaN

4.3µm

5.0µm

5.0µm
(c)

SiC

5.0µm

regrown GaN

nanocolumns
5.0µm

SiO2
Fig. 4. Top SEM images at time of regrowth of GaN layer on nanocolumn

substrate: (a) before growth, (b) during growth (10 min growth), and(c) after growth (180 min growth).

AlN
SiC

SiO2 mask. Accordingly, we found that the deep-throughhole-patterned mask used in this experiment is eﬀective forgrowing uniform nanocolumn crystals.
Then, a planarized GaN layer was regrown on thesubstrate with GaN nanocolumns, after etching the SiO2
mask down to 100 nm thickness. Figure 4 shows the
regrowth process of this sample. From (a) to (b), islands
of GaN are randomly growing to form {1100}
and {1101}
facets from the side walls of the nanocolumns. Island
crystals are formed in accordance with the variation in Ga
diﬀusion length. Because the lateral growth is dominant inthis situation, the height of the island crystals is almost kept
constant at 500 nm from the mask level. From Fig. 4, theplanarization of the regrown GaN is conducted by the
coalescence of the island crystals growing in the lateral
direction. A cross-sectional SEM image of the GaN layer isshown in Fig. 5. The thickness of the regrown GaN layer
was 4.3 m. A ﬂat surface and the absence of void inclusions
around the regrowth interface were observed.
Finally, CL mapping images of the GaN layers with andwithout GaN nanocolumns are shown in Fig. 6. The dark
spots observed in the images correspond to threading
dislocations. From the images, the dislocation density of
the GaN layer directly grown on the ﬂat AlN/SiC substrate
was 9:3  108 cm 2 , and that of the GaN layer grown on

GaN layer

1.0µm
Fig. 5. Cross-sectional SEM image of regrown GaN layer with thicknessof 4.3 m and magniﬁed view around nanocolumn interface.

the GaN nanocolumn template was 1:3  108 cm 2 . A
signiﬁcant reduction in dislocation density was realized bynano-ELO.
The dislocation density of a commercial SiC substrateis about 1:0  105 cm 2 . However, that of the AlN buﬀer
layer grown on the SiC substrate is about 1010 cm 2 , owing
to lattice mismatch. The number of dislocations decreases
when a GaN layer is grown on the AlN buﬀer layer.
The reduction in dislocation density is thought to be dueto the fact that nano-ELO prevents the propagation ofinternal dislocations existing in the AlN buﬀer layer.
Furthermore, the GaN layer regrown by a lateral growthover the SiO2 mask also has a small number of dislocations,similarly to the conventional ELO-grown GaN layer. The
reduction in the number of dislocations in the overgrown
GaN layer is partly thought to be caused by the initial GaN
nanocolumn crystals with a low dislocation density. In the
regrowth stage, island crystals are formed uniformly. Theseisland crystals coalesce with the termination of dislocations

on AlN/SiC substrateon GaN nanocolumn template

CL images

5µm

5µm

Dislocation density [cm-2]

9.4 × 10 8

1.3 × 10 8

Fig. 6. CL mapping images of GaN layer on ﬂat AlN/SiC substrate and GaN nanocolumn template.

08JE23-3# 2013 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 52 (2013) 08JE23

S. Umeda et al.

and are eventually planarized, similarly to those in theconventional ELO with a micron scale pattern.19–21) With the
optimization of the conditions for the regrowth of GaN,
further reduction in dislocation density can be expected.

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4. Conclusions

231104.

We demonstrate a new GaN nanocolumn growth method
using deep SiO2 through-holes. This gives reproducible anduniform GaN nanocolumn crystals under typical MOVPE
growth conditions. By applying this method, the GaN layer
regrown on the nanocolumns shows a low dislocation
density of 1:3  108 cm 2 , which is much lower than that
9:3  108 cm 2 of the GaN layer grown on a ﬂat AlN/SiC
substrate.

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# 2013 The Japan Society of Applied Physics

pssPhys. Status Solidi RRL 6, No. 4, 145–147 (2012) / DOI 10.1002/pssr.201206038

Properties of nitride-based
photovoltaic cells under concentrated
light illuminationwww.pss-rapid.com

Shota Yamamoto1, Mikiko Mori1, Yosuke Kuwahara1, Takahiro Fujii1, Tatsuo Nakao1, Shinichiro Kondo1,
Motoaki Iwaya*, 1, Tetsuya Takeuchi1, Satoshi Kamiyama1, Isamu Akasaki1, 3, and Hiroshi Amano2, 3
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received 31 January 2012, revised 7 March 2012, accepted 7 March 2012
Published online 13 March 2012
Keywords nitride semiconductors, solar cells, concentrator cells, GaInN*

Corresponding author: e-mail iwaya@meijo-u.ac.jp

We investigated the properties of nitride-based solar cells under concentrated light illumination from 1 to 200 suns. The
conversion efficiency of our solar cells increased with increasing concentration up to 200 suns. The short-circuit cur-rent density, open-circuit voltage, fill factor, and conversion
efficiency were 510 mA/cm2, 1.9 V, 70%, and 3.4%, respectively, under an air mass filter of 1.5G at 200 suns and room
temperature.

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 General Since the band gap of GaInN ternary alloys
covers a broad range from 0.65 to 3.43 eV [1], these alloys
are suitable for solar cell applications. So far, we have succeeded in fabricating GaInN-based solar cells [2, 3]. By
reducing the pit density in the GaInN active layer using
freestanding c-plane GaN substrates and applying GaInN
superlattice structures, the conversion efficiency of our nitride-based solar cells has reached 2.9% [4].
One of the disadvantages of compound semiconductor
solar cells is their high cost compared with that of Si or organic solar cells. This problem can be solved by using acondenser lens system. In general, a low-cost material canbe used for the condenser lens system. Thus, if the solar
cells can operate at 200 suns, for example, the cost per chip
would be reduced almost 200-fold. In addition, the performances of solar cells using Si, InP, and GaAs are improved by focusing the sunlight [5], although the improvement in performance is limited depending on a material, a junction quality, a cell configuration, and an electrode pattern. The maximum efficiencies of Si and InP solar cells have been reported to be obtained at 90–95 suns[5]. In contrast, the maximum efficiency of GaAs solar
cells has been reported to be obtained at over 200 suns [5].
On the other hand, photovoltaic cell characteristics of

GaInN-based solar cells have been investigated in concentration up to 30 suns only [6]. Details of nitride-based solar
cell characteristics in concentration above 30 suns are still
unknown.
In this study, we investigate the focusing properties ofnitride-based solar cells up to 200 suns. We also discuss
the dependence of the concentration ratio on the solar cellcharacteristics.
2 Sample structure Nitride-based solar cells weregrown by metal organic vapour phase epitaxy. Trimethylindium, trimethylaluminium, trimethylgallium, triethylgallium, and ammonia were used as the source gases.
Figure 1 schematically shows the structure of the devices prepared in this study. We grew 50 pairs of unintentionally doped Ga0.83In0.17N (3 nm)/Ga0.93In0.07N (0.6 nm)
superlattice layers as active layers on freestanding GaN
substrates. Another 10 pairs of Si-doped Ga0.90In0.10N(3 nm)/GaN (3 nm) superlattice layers were inserted beneath the active layers. The Si concentration in the 10 pairsof superlattice layers was 3 × 1018 cm–3. The role of theadditional Si-doped superlattice layers was to reduce thedislocation density and to obtain high-quality active layerswith a higher InN mole fraction while maintaining a low© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica146

rrl

S. Yamamoto et al.: Nitride-based photovoltaic cells under concentrated light illumination

Figure 1 (online colour at: www.pss-rapid.com) Schematic viewof sample structure of nitride-based solar cells.

series resistance (RS) [3]. In addition, the numbers of superlattice layers are experimentally obtained. Dislocation
density was reduced also by using the GaN substrate. A
semitransparent Ni (5 nm)/Au (5 nm) ohmic contact and aTi (30 nm)/Al (100 nm)/Ti (20 nm)/Au (150 nm) ohmic
contact were formed on p-GaN (the surface of the device)and on n-GaN (the reverse side of the device) by electron
beam evaporation, respectively. The devices had a verticalstructure.
3 Experimental conditions The conversion efficiencies of the devices were measured using a solar simulator with an air mass (AM) filter of 1.5G (Asahi Spectra
HAL-302). In this solar simulator, the infrared light waseliminated with a filter. Both the irradiation area and thephoton density were changed to adjust the concentration of
the light from 1 to 200 suns. The current density versus
voltage characteristics of the devices were measured from
1 to 200 suns at room temperature (RT). No active temperature control of the devices was performed.
4 Experimental result Figure 2 shows the current
density versus voltage characteristics of the nitride-basedsolar cells under the solar simulator (1 to 200 suns) at RT.
We found that the open-circuit voltage (VOC) of the nitridebased solar cells was increased with increasing concentra-

Open-circuit voltage[V]

2.5
2.0
1.5
1.0
0.5
0.0

VOC
VOCfrom Eq.(1)

1

10

100

Concentration ratio
Figure 3 (online colour at: www.pss-rapid.com) VOC as functionof concentration ratio. The red dashed line was obtained using
Eq. (1).

tion ratio and appeared to saturate above 50 suns. VOC isplotted as a function of concentration ratio in Fig. 3.
In general, VOC can be expressed as [7]VOC =nkT Ê JSC CR ˆln Á
+ 1˜ ,
Ë J0
¯
q

(1)

where JSC, CR, J0, n, k, T, and q are the short-circuit current density, concentration ratio, reverse saturation current,
n-value (a correction value of the diode equation), Boltzmann constant, temperature, and elementary charge, respectively. Equation (1) shows that VOC should increase
logarithmically with increasing light intensity. The dashedline in Fig. 3 almost matches the experimental results for 1
to 200 suns.
Figure 4 shows the fill factor (FF) and JSC as functionsof the concentration ratio. FF only changed slightly, from72 to 70%, when the concentration was increased from 1 to
200 suns. In general, FF under concentrated sunlight is affected by the enhanced carrier recombination in the inter1.0

600
JSC
FF

0.8

400

0.6

300
0.4

200

0.2

100
0

Figure 2 (online colour at: www.pss-rapid.com) Current density
versus voltage characteristics of nitride-based solar cells under
the solar simulator (AM of 1.5G, 1 to 200 suns) at RT.
© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

FF

JSC[mA/cm2]

500

0

50

100

150

0.0
200

Concentration ratio
Figure 4 (online colour at: www.pss-rapid.com) JSC and FF as
functions of concentration ratio.
www.pss-rapid.com

RapidResearch Letter
Phys. Status Solidi RRL 6, No. 4 (2012)

147

Table 1 Device performance characteristics of nitride-basedsolar cells under the solar simulator (AM of 1.5G, 1 and 200suns).

1 sun
200 suns

VOC
[V]

FF
[%]

JSC
η
[mA/cm2] [%]RS
RSH[Ω cm2] [Ω cm2]

1.5
1.9

72
70

1.4
520

110
0.53

1.5
3.4

12000
18

face region due to the high carrier density under concentrated sunlight [7].
In Ref. [6], FF decreased from 64% to 57% when theconcentration ratio increased from 1 to 30 suns, whereas
for the samples used in this study, FF did not decrease significantly when the concentration ratio increased. Previous
reports concluded that the significant decrease in FF withincreasing concentration was due to the poor crystal quality of the sample [6]. Our samples with the GaN substrateand superlattice structure have high crystallinity [3].
The dislocation density of this sample was approximately
5 × 107 cm–2. Therefore, a high FF was maintained up to
200 suns.

JSC was approximately proportional to the concentration ratio and increased 370-fold when the concentration
ratio was increased from 1 to 200 suns (cf. Table 1).
Figure 5 shows the measured series resistance (RS) andshunt resistance (RSH) as functions of the concentration
ratio. We calculated RS and RSH from the slope of theI–V curves [8]. RS and RSH rapidly decrease with increasing
concentration ratio. Because the photo current of this solar
cell is very small at low sun, RS and RSH appear larger. Thus,
RS and RSH are rapidly reduced due to an increase in thephoto current with high concentrated light illumination.
Figure 6 shows the power conversion efficiency η as afunction of concentration ratio. We also summarize thevalues of JSC, VOC, RS, RSH, FF, and η at 1 and 200 suns exposures in Table 1. The efficiency increases from 1.5% to3.4% when the concentration ratio increases from 1 to
200 suns. The efficiency is increased by a factor of 2.3,
whereas VOC is increased by a factor of 1.3. The higher rateof increase of the ratio in η than that in VOC is due to theincrease in JSC with increasing concentration ratio. That is,
the nitride-based solar cells operate effectively even at
200 suns. The conversion efficiency of 3.4% is the highest
ever reported value for nitride solar cells.
5 Summary In conclusion, we investigated the nitride-based solar cells under concentrator conditions. VOC,
JSC, and η monotonically increased from 1 to 200 suns,
while FF decreased slightly. JSC, VOC, FF, and η were
510 mA/cm2, 1.9 V, 70%, and 3.4%, respectively, at AM
1.5G and 200 suns at RT.
Acknowledgement This study was partially supported bya New Energy and Industrial Technology Development Organization (NEDO) Project on the Research and Development of Innovative Solar Cells.

References

Figure 5 (online colour at: www.pss-rapid.com) RS and RSH asfunctions of concentration ratio (1 to 200 suns).

Figure 6 η as a function of concentration ratio.
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E. E. Haller, H. Lu, W. J. Schaff, W. K. Metzger, andS. Kurtz, J. Appl. Phys. 94, 6477 (2003).
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[5] M. A. Green, K. Emery, D. L. King, Y. Hishikawa, andW. Warta, Prog. Photovolt.: Res. Appl. 14, 45 (2006).
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© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

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Applied Physics Express 10, 025502 (2017)https://doi.org/10.7567/APEX.10.025502

Relationship between lattice relaxation and electrical properties in polarization dopingof graded AlGaN with high AlN mole fraction on AlGaN template
Toshiki Yasuda1, Tetsuya Takeuchi1*, Motoaki Iwaya1, Satoshi Kamiyama1, Isamu Akasaki1,2, and Hiroshi Amano2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8062, Japan
3
Institute of Materials and Systems for Sustainability, Nagoya University, Nagoya 464-8062, Japan2*E-mail: take@meijo-u.ac.jp
Received November 17, 2016; accepted January 11, 2017; published online January 30, 2017
We investigated polarization doping for hole generation in abrupt and graded GaN/Al0.7Ga0.3N interfaces on Al0.99Ga0.01N templates. The abrupt
interface exhibited hole generation, whereas the graded interface exhibited electron generation. In the graded AlxGa1%xN (x = 0.65–0), a gradedpart with an AlN mole fraction ranging from 0.2 to 0 showed a large relaxation. Theoretical estimation revealed that this part contained positive
polarization charges, accumulating electrons. Via Mg doping in the graded AlGaN layer, we obtained a high hole concentration of 3 ' 1013 cm%2.
These results indicate that understanding the relaxation conditions in the graded layer is indispensable for polarization doping.
© 2017 The Japan Society of Applied Physics

itride semiconductor-based deep-ultraviolet (DUV)
light-emitting diodes (LEDs) have been proposed toreplace UV lightbulbs. However, the external quantum eﬃciency of UV LEDs is considerably low, and it is more
pronounced in the DUV region.1–3) Among the reasons forthe low eﬃciency are the poor electrical properties of p-type
AlGaN with a high AlN mole fraction. Nitride semiconductors can exhibit p-type conductivity after Be, Zn, or Mg
doping.4–7) In general, Mg is mostly used to form p-type
AlGaN.8,9) However, it is almost impossible to generateholes in AlGaN layers with high AlN mole fractions, because
acceptors such as Mg in such AlGaN layers have large
ionization energies.10,11) For instance, the resistivity of
Mg-doped Al0.7Ga0.3N exceeds 1,000 Ω·cm. This is why
most nitride DUV LEDs contain a relatively thick p-type
GaN, providing holes in the active regions and an ohmic
contact to the p-electrode.12–18) However, GaN absorbs UVlight with wavelengths shorter than 365 nm; thus, there isa tradeoﬀ between the electrical and optical properties.
Polarization doping has been proposed as an alternative
approach for generating carriers in nitride semiconductors.19)
The polarization charge concentration, NP_2D (cm−2), is estimated as NP_2D = P=q, where P (C=cm2) is the polarization
charge density, and q (C) is the elementary charge. The polarization charge density includes spontaneous and piezoelectric
polarization components.20,21) The polarization charge concentration accumulates the free carrier concentration, N2D, with
opposite signs to satisfy the charge neutrality, N2D ≅ NP_2D. For
instance, two-dimensional (2D) electron gas has been accumulated to the polarization positive charges at an AlGaN=GaN
interface in a heterostructure ﬁeld-eﬀect transistor. Furthermore, 2D hole gas was observed at the other GaN=AlGaN
interface. In addition, Simon et al. proposed Mg-doped graded
AlxGa1−xN (x = 0.3–0) structures in which holes are threedimensionally distributed, so that the polarization doping is also
useful in optoelectronic devices that require vertical conduction.22) For the graded structure, the three-dimensional (3D)
polarization charge concentration, NP_3D (cm−3), is estimated asN

NP 3D ¼@P;
q  @dwhere d (cm) is the thickness of the graded layer.19) Thus far,
we have shown that the hole concentrations in polarization

doping can be tuned by controlling the AlN mole fractions
in GaN=AlGaN structures up to 0.34 AlN.23) Polarization
dilution involves compensating unnecessary polarization
charges (e.g., positive charges in the p-type layer) with asuﬃcient amount of ionized impurities (e.g., ionized acceptors) for improving the vertical conductivity.24) These results
have been obtained with AlGaN on GaN templates.
Recently, polarization doping on AlN templates, which
are more suitable than GaN templates for DUV LEDs, hasbeen reported. Zhang et al. demonstrated hole generationin Mg-doped graded AlxGa1−xN (x = 0.3–0).25) Li et al.
obtained a high hole concentration in Be-doped graded
AlxGa1−xN (x = 1–0.7).26) However, p-type graded AlGaN
with a substantially wider AlN mole fraction range, e.g.,
0.7 to 0, should be investigated with practical DUV-LEDs
to simultaneously obtain a high hole concentration and a
low contact resistivity for the p-electrode. In this study, weinvestigated polarization doping in GaN=Al0.7Ga0.3N structures on AlN templates. We measured hole accumulations
to the polarization charges in abrupt and graded GaN=
Al0.7Ga0.3N interfaces with and without Mg doping withrespect to the measurement temperature. We found a lattice
relaxation in the graded AlGaN layer with AlN mole fractions ranging from 0.2 to 0 and demonstrated that this lattice
relaxation aﬀected the conductivity type.
All samples were prepared on AlN templates grown on
c-plane sapphire substrates by metalorganic vapor phase
epitaxy. The AlN templates were grown with the addition ofa small amount of trimethylgallium to improve the surface
morphology, resulting in a 2-µm Al0.99Ga0.01N layer on an AlN
buﬀer layer.27) We then prepared three diﬀerent Hall samples.
The ﬁrst had an abrupt interface comprising a 100-nm layer of
undoped Al0.74Ga0.26N, a 10-nm layer of undoped GaN, and a
40-nm layer of p-GaN on the Al0.99Ga0.01N template, as shownin Fig. 1(a). The second Hall sample had a graded interface
comprising a 100-nm layer of undoped Al0.65Ga0.35N, a 100-nm
layer of undoped graded AlxGa1−xN (x = 0.65–0), and a 40-nm
layer of p-GaN on the Al0.99Ga0.01N template, as shown in
Fig. 1(b). The third Hall sample was similar to the second one
but contained a 100-nm layer of Mg-doped (2 × 1019 cm−3)
graded AlxGa1−xN (x = 0.65–0) instead of the undoped
AlxGa1−xN. The 40-nm p-GaN layers were required as
p-contact layers to inject the current. For X-ray diﬀraction

025502-1© 2017 The Japan Society of Applied Physics

Appl. Phys. Express 10, 025502 (2017)(a)T. Yasuda et al.

(b)

(a)

(b)

Fig. 1. Hall-sample structures containing (a) the abrupt GaN=
Al0.74Ga0.26N heterostructure and (b) the undoped and Mg-doped graded
AlxGa1−xN (x = 0.65–0) layers.

(XRD) measurements, another sample was prepared on theAl0.99Ga0.01N template, comprising a 100-nm layer of undoped
Al0.65Ga0.35N and a 100-nm layer of Mg-doped graded
AlxGa1−xN (x = 0.65–0.3) without a p-GaN contact layer.
Hall-eﬀect measurements were performed to determine thecarrier concentrations with respect to the temperature. Asymmetric reciprocal space mappings (RSMs) were obtained via
XRD. The AlN mole fractions in the graded AlGaN layer
were estimated using secondary ion mass spectrometry (SIMS).
Figure 2 shows the carrier concentrations of the three
Hall samples with respect to the temperature. As shown in
Fig. 2(a), the ﬁrst Hall sample with the abrupt interface
showed p-type conductivity (with red dots) throughout thetemperature range. The sheet hole concentration at room
temperature was 1.5 × 1013 cm−2, showing almost no temperature dependence.23) For comparison, the sheet carrier concentrations of a 100-nm Mg-doped p-GaN layer are plottedwith black dots in the same ﬁgure, showing a signiﬁcanttemperature dependence due to the large ionization energy of
Mg. Thus, we conclude that the observed holes were accumulated to the polarization negative charges at the GaN=Al0.74Ga0.26N interface. On the other hand, as shown in Fig. 2(b),the second Hall sample with the undoped graded AlGaN
unexpectedly showed n-type conductivity (with blue dots)with a weak temperature dependence. The sheet electron concentrations were ∼1 × 1013 cm−2 in the measured temperature
range. As shown in Fig. 2(c), the third Hall sample with
Mg-doped graded AlGaN exhibited p-type conductivity witha sheet hole concentration of 3 × 1013 cm−2 at room temperature. The 3 × 1013 cm−2 sheet hole concentration corresponds to a volume hole concentration of 3 × 1018 cm−3,
considering the graded 100-nm distribution. The conductivity
type changed from p-type (with red dots) to n-type (withblue dots) as temperature was reduced below 240 K.
The structural properties of the Hall samples were investigated to investigate the reason for this conductivity-type
behavior. First, polarity inversion was suspected; thus, we
performed KOH wet etching (80 °C, 1 min) on all the Hall
samples. No changes in the surface morphologies of thesamples were observed, suggesting that the samples had
typical Al polarity and that the polarity inversion was not acause of the n-type conductivity. Second, the AlN mole
fractions in the Mg-doped graded AlGaN layer of the third
Hall sample were measured via SIMS. Figure 3 shows that
the AlN mole fractions decreased monotonically towards the
Al-polar surface side, as designed, indicating that no net
positive polarization charges were generated and thus no(c)Fig. 2. Sheet carrier concentrations with respect to the temperature for(a) the abrupt GaN=AlGaN structure, (b) the undoped graded AlGaN layer,and (c) the Mg-doped graded AlGaN layer. Red and blue dots indicate thehole and electron concentrations, respectively.

electrons could be accumulated. Third, the strain and relaxation conditions of the lattices in the samples were investigated.
 XRD RSM of the ﬁrst Hall
Figure 4(a) shows the ð2024Þ
sample with the abrupt GaN=Al0.74Ga0.26N interface. The
results clearly indicate that the underlying Al0.74Ga0.26N was
5% relaxed and that the GaN was fully relaxed. The theoreticalpolarization charge concentrations were calculated to be
3.7 × 1013 and 1.7 × 1013 cm−2 for the fully strained GaN(to the Al0.74Ga0.26N) and fully relaxed GaN, respectively.
The measured hole concentration of 1.5 × 1013 cm−2 shownin Fig. 2(a) is well-explained by the value (1.7 × 1013 cm−2)
for the fully relaxed GaN. Figures 4(b) and 4(c) show thecorresponding RSMs for the third Hall sample with the
Mg-doped graded AlxGa1−xN (x: 0.65–0) layer and the other
XRD sample with the Mg-doped graded AlxGa1−xN (x =0.65–0.3) layer. Figure 4(c) indicates that the underlying
Al0.65Ga0.35N was 20% relaxed and that the graded AlGaN
with AlN mole fractions ranging from 0.65 to 0.3 was coherently grown on the underlying Al0.65Ga0.35N. In contrast,
Fig. 4(b) suggests that the graded AlGaN with AlN mole
fractions ranging from 0.65 to 0 comprised the following
parts: 1) the graded AlGaN with an AlN mole fraction ranging
from 0.65 to 0.2, which was fully strained to the Al0.65Ga0.35N
layer, and 2) the graded AlGaN with AlN mole fractions
ranging from 0.2 to 0, which was gradually and then fully
relaxed with increasing proximity to the GaN at the surface.
According to the strain and relaxation conditions observedin the graded AlGaN layer, the theoretical absolute polarization charge concentrations (cm−2) and net polarization
charge concentrations (cm−3), NP_3D , of the graded AlGaN
layer were calculated with respect to the depth (or AlN mole
fraction) of the graded AlGaN layer, as indicated by the solidlines in Figs. 5(a) and 5(b), respectively. For comparison, the025502-2© 2017 The Japan Society of Applied Physics

Appl. Phys. Express 10, 025502 (2017)

T. Yasuda et al.
(a)Fig. 3. AlN mole-fraction proﬁle for the Mg-doped graded AlGaN layer,
measured by SIMS.

(b)

(a)

(b)
Fig. 5. (a) Theoretical absolute polarization charge concentrations (cm−2)and (b) theoretical measurable net polarization charge concentrations (cm−3)
for the partially relaxed graded AlGaN layer.

Finally, we theoretically estimated the carrier concentration of the Mg-doped graded AlGaN with an AlN mole fraction ranging from 0.2 to 0, which was gradually relaxed.
Because ionized Mg acceptors — rather than electrons —should neutralize the positive polarization charges,24) wesolved the following charge-neutrality condition using
Shockley’s graphical method:28)

(c)

þ
þ
n þ N
A ¼ p þ ND þ NP ;EC  EFn ¼ NC exp ;kT
EF  EV
p ¼ NV exp ;kT XRD RSM patterns of (a) the abrupt GaN=AlGaN
Fig. 4. ð2024Þ
structure, (b) the Mg-doped graded AlxGa1−xN layer (x = 0.65–0), and (c) the
Mg-doped graded AlxGa1−xN layer (x = 0.65–0.3).

theoretical absolute polarizations in the cases of the AlGaN
fully strained to the Al0.65Ga0.35N layer and the fully relaxed
AlGaN are plotted in Fig. 5(a) with dashed and dotted lines,
respectively. As shown in Fig. 5(a), because the part with
AlN mole fractions ranging from 0.65 to 0.2 was fullystrained to the Al0.65Ga0.35N layer, the values in this part are
the same as those in the strained case. In contrast, the partwith the AlN mole fraction ranging from 0.2 to 0 is graduallyand then fully relaxed so that the absolute polarization charge
concentrations in this part gradually change from the valueof the strained Al0.2Ga0.8N to the value of the relaxed GaN,
resulting in the opposite sign of the slope compared with thesimply strained and simply relaxed cases. The aforementioned 3D polarization charge concentration NP_3D containing
the ∂P=∂d component is plotted in Fig. 5(b). This ﬁgure
reveals that the net polarization charge concentration in thegradually relaxed part (graded AlGaN with AlN mole fractions ranging from 0.2 to 0) has the opposite sign (positive)
with a polarization charge concentration of ∼3.4 × 1018 cm−3.
This suggests that electrons — no longer holes — were accumulated on the graded AlGaN with AlN mole fractions
ranging from 0.2 to 0. The measurable electron sheet concentration was theoretically estimated to be 1 × 1013 cm−2 byconsidering the thickness of the graded AlGaN (3.4 × 1018
cm−3 × 30 nm). This value agrees well with the measured
electron value (1 × 1013 cm−2), as shown in Fig. 2(b).

Nþ
D ¼

N
A ¼

ND

;
EF  ED1 þ 2 exp
kT
NA

:
EA  EF1 þ 4 exp
kT

ð1Þ
ð2Þ
ð3Þ
ð4Þ

ð5Þ

Equation (1) expresses the charge-neutrality condition, including the positive polarization charge. Here, n and p are theelectron and hole concentrations, respectively; ND and NA are
the donor and acceptor concentrations (0 and 2 × 1019 cm−3),
respectively; N þD and N A are the ionized-donor and ionizedacceptor concentrations, respectively; N þP is the positive
polarization charge density (3.4 × 1018 cm−3); NC and NV are
the eﬀective density of states at the conduction and valencebands, respectively; EA and ED are the acceptor and donor
energy levels (0.2 and 3.58 eV), respectively; EV and EC are thevalence- and conduction-band edge energy levels (0 and 3.6eV), respectively; k is Boltzmann’s constant, and T is the temperature. For simplicity, we used the average parameters for
the Mg-doped graded AlGaN layer, i.e., Al0.1Ga0.9N. Figure 6
shows the total negative [left side of Eq. (1): blue line] andpositive [right side of Eq. (1): red line] charge densities withrespect to the Fermi energy in the Mg-doped graded AlGaN.

025502-3© 2017 The Japan Society of Applied Physics

Appl. Phys. Express 10, 025502 (2017)

T. Yasuda et al.

(a)layers, which explains the observed electron generation in the
undoped graded AlGaN layer with a wide grading range of
AlN mole fractions. A combination of Mg doping and polarization doping resulted in a high sheet carrier concentration of
3 × 1013 cm−2 (or 3 × 1018 cm−3 in 100 nm), which is usefulfor suﬃcient hole injections in DUV LEDs.
Acknowledgments This study was partially supported by the MEXT
Program for the Strategic Research Foundation at Private Universities(2012–2016), MEXT-Grants-in-Aid for Scientiﬁc Research and Specially
Promoted Research (No. 25000011), JSPS Grant-in-Aid for Scientiﬁc Research A(No. 15H02019), JSPS Grant-in-Aid for Scientiﬁc Research B (No. 26286045),
JSPS KAKENHI (No. 16H06416), JST CREST (No. 16815710), and the Nippon
Sheet Glass Foundation for Materials Science and Engineering.

(b)

Fig. 6. Calculated total negative=positive charges and components of thepositive charges in the graded AlGaN with a mole fraction ranging from 0.2
to 0 with respect to the Fermi energy at (a) 300 K and (b) 200 K. The energy
the valence-band edge is 0 eV.

The intersection satisﬁes the charge-neutrality condition for agiven Fermi energy. Figure 6 also shows the two componentsof the total positive charges, the hole concentrations (dottedblack line), and the positive polarization charge concentrations(dashed black line). At room temperature, the Fermi energy of
0.198 eV yielded the charge-neutrality condition, resulting in
p-type AlGaN with a hole concentration of 1.4 × 1016 cm−3, asshown in Fig. 6(a). Considering that the hole concentration of
the strained graded AlGaN with an AlN mole fraction ranging
from 0.65 to 0.2 was 3.4 × 1018 cm−3, we calculated the measurable hole concentration of the third Hall sample to be
2.4 × 1013 cm−2 (3.4 × 1018 cm−3 × 70 nm + 1.4 × 1016 cm−3 ×
30 nm), which agrees well with the measured value (3 ×1013 cm−2) shown in Fig. 3(c).
To elucidate the reason for the observed change in the conductivity from p- to n-type around 240 K, Fig. 6(b) describes
the charge-neutrality condition for the same case at 200 K.
The hole concentration clearly decreased to 2 × 1014 cm−3.
However, the Fermi energy remained relatively constant,indicating p-type conductivity. Therefore, the observed conductivity switching in the third Hall sample cannot presently
be explained. Furthermore, the reason why the observed
electron concentration of the third Hall sample at a lowtemperature was higher than that of the second Hall sample
is unknown. Nevertheless, we point out that the strain andrelaxation conditions drastically aﬀect not only the magnitude but also the type of conductivity in the graded layers
for polarization doping. To properly utilize the polarization
doping in the graded layers, we must further investigate thestrain and relaxation conditions.
In this study, we investigated polarization doping for hole
generation in GaN=Al0.7Ga0.3N structures on AlN templates.
The abrupt GaN=Al0.7Ga0.3N interface generated holes,
whereas the graded GaN=Al0.7Ga0.3N interface generated
electrons. The lattice relaxations drastically aﬀected not only
the magnitude but also the type of conductivity in the graded

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025502-4© 2017 The Japan Society of Applied Physics

1

1

solidiawww.pss-a.com

applications and materials science

Polarization dilution in a Ga-polar
UV-LED to reduce the inﬂuence ofpolarization charges,1

statusphysicaPhys. Status Solidi A 212, No. 5, 920–924 (2015) / DOI 10.1002/pssa.201431730

1

1

Toshiki Yasuda* , Kento Hayashi , Syouta Katsuno , Tetsuya Takeuchi , Satoshi Kamiyama ,Motoaki Iwaya1, Isamu Akasaki**,1,2, and Hiroshi Amano**,2,3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8062, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8062, Japan
2

Received 30 September 2014, revised 20 November 2014, accepted 25 November 2014
Published online 15 January 2015
Keywords AlGaN, light-emitting diodes, polarization* Corresponding author: e-mail 143441506@ccalumni.meijo-u.ac.jp, Phone: þ81-52-832-1151**

Isamu Akasaki and Hiroshi Amano are laureates of the Nobel Prize in Physics 2014.

We investigated the inﬂuence of polarization charges innitride-based semiconductors. The inﬂuence due to polarization charges was calculated excluding the inﬂuences of theband offset. We found that the polarization charges (1  1013
cm2) resulted in an energy spike of more than 100 meV at thelocation of the charges, which is a similar value to the band
offset. We then proposed the concept of polarization dilution to
suppress the energy spike for better hole transport by using a

graded Mg-doped AlGaN layer in UV-LEDs. Device
simulation results indicate lower operating voltage and higher
injection efﬁciency by using the polarization dilution. So far,our actual 350 nm LED with the polarization dilution showed
lower operating voltage. These results suggested such polarization-charge management is important in the design of thenitride semiconductors.

ß 2015 WILEYVCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction AlGaN-based ultraviolet (UV) lightemitting diodes (LEDs) are expected to be highly efﬁcient,reliable, and compact light sources for sterilization, water
puriﬁcation, printing, and so on. Currently, external quantum
efﬁciencies (EQEs) of the UV-LEDs are much lower thanthose of commercialized visible LEDs. For instance, thereported highest EQEs of UV-A, UV-B, and UV-C are 44% at
365 nm [1], 10.4% at 278 nm [2], 0.2% at 227 nm [3],respectively. One of the reasons for such low EQEs is theinferior p-type characteristic of AlGaN resulting in poor hole
injection into the active region. While the ionization energy of
Mg acceptors in GaN is 170 meV [4], that in AlN is
600 meV [5]. Therefore, p-type AlGaN especially with high
AlN mole fraction shows very low hole concentrations [6].
On the other hand, the nitride semiconductors are wellknown to show large polarization characteristics. Large
spontaneous [7] and piezoelectric [8] polarization charges
are generated, and remain at heterointerfaces. As these charges
are ﬁxed charges, they do not contribute as mobile carriers inprinciple. The concentrations of the polarization charges in the

nitride semiconductors, however, are comparable to or evenhigher than those of electrons and holes in impurity-doped
layers. Thus, the polarization charges largely inﬂuence themobile carriers as well as energy-band structures in nitride
semiconductors. A favorable example to utilize the polarization
charges is a two-dimensional electron gas (2DEG) in aheteroﬁeld effect transistor (HFET) [9]. The 2DEG isaccumulated even without any impurity doping at one of the
AlGaN/GaN interfaces where positive polarization charges are
generated. For instance, as the concentration of the polarization
charges at the Al0.2Ga0.8N/GaN interface is more than1  1013 cm2, almost the same high number of the electrons
is built up. Furthermore, the polarization charges with the
opposite sign are also generated at the other side of the interface,and now 2D hole gas is accumulated there. In addition, a newconcept, “polarization doping” [10] was proposed, in whichsuch a large number of the polarization “sheet” charges are
three-dimensionally distributed by forming compositional
graded interfaces, resulting in a three-dimensional (3D) carrier
distribution. With the polarization doping, a compositionallyß 2015 WILEYVCH Verlag GmbH & Co. KGaA, Weinheim

Original
Paper
Phys. Status Solidi A 212, No. 5 (2015)

921[0001]

graded-AlGaN
P-Al0.2Ga0.8N

increase
Al composion

Posive charges Negave charges= 1x1013cm-2= 1x1013cm-2

Figure 1 Assumed structure (putting positive and negative
charges in p-Al0.2Ga0.8N).

Polarizaondipolar

Polarizaon(+)
charges

unit cell(-)

graded p-AlGaN showed a much higher 3D hole concentration
allowing vertical carrier transport [11]. Then, this concept wasproved in a light emitting device, an N-polar UV-LED, showing
higher hole injection [12].
In our previous study, we attempted to use polarization
doping in a more conventional Ga-polar blue LED [13].
Even though we clearly observed high hole accumulation at
an AlGaN graded layer with a separate Hall measurement,
the blue LED with the similar graded layer, whichcorresponded to an electron blocking layer, showed very
low output power. This fact indicates that the polarization
doping works in N-polar LEDs but not in Ga-polar LEDs.
The difference between N-polar and Ga-polar in terms of thepolarization is the sign of the polarization charges. The
polarization charges with the opposite sign are generated at
the equivalent interfaces in the same device structure.
In this study, we ﬁrst calculated the inﬂuence of thepositive and negative polarization charges on energy-band
proﬁles in the impurity-doped layer. Then, we proposed atechnique to reduce the inﬂuence of the polarization
charges, “polarization dilution”. Finally, we performed
theoretical and experimental investigations on the effect ofthe polarization dilution in 350 nm LED.
2 Inﬂuence of polarization First, we calculated theinﬂuence of polarization charges with a commercially available
device simulator (SiLENSe) implementing a 1D drift-diffusion
model in consideration of the strong piezoelectric/spontaneouspolarization charges. Here, we are interested in just the inﬂuenceof the polarization charges on the energy-band proﬁles. Thus, weNetpolarizaon(+)
charges(-)
All ﬁxed charges(+)(in p-typegraded-AlGaN)

Posive polarizaon charges(-)

Ionized acceptors

Figure 3 Mechanism of polarization dilution. Positive polarization charges are distributed in a graded layer by grading withincrease of Al composition. Positive polarization charges arecompensated by ionized acceptors under charge-neutrality
conditions.

assumed a simple uniform p-type material, Mg-doped
Al0.2Ga0.8N, containing positive and negative sheet charges,
1  1013 cm2, separately located, as shown in Fig. 1. The
assumed structure allows us to solely estimate the inﬂuence of thecharges, excluding that of a band offset that must be formed at theheterostructure interface. The ﬁxed charges, 1  1013 cm2,
correspond to the number of polarization charges at theAl0.2Ga0.8N/Al0.4Ga0.6N interface.
The calculated valence-band energy proﬁle of the assumed
structure is plotted in Fig. 2. An energy spike and an energy dip
to the holes in the valence band are observed at the locations of

p-GaN(30nm)(a)p-Al0.4Ga0.6N(20nm)

p-Al0.2Ga0.8N (100nm)

Graded-p-AlxGa1-xN(x:0.2→0.4,20nm)

Electron blocking layer
u-Al0.2Ga0.8N (last barrier layer)

0.2 0.4

GaN/n-Al0.2Ga0.8N 5QWs

n-Al0.2Ga0.8N(b)p-Al0.4Ga0.6N(20nm)

AlN
0.2 0.4
Sapphire

Figure 2 The calculated valence-band energy proﬁle of theassumed structure (Fig. 1). Energy spike and dip for hole is
160 meV and 140 meV, respectively.
www.pss-a.com

Al composion proﬁleFigure 4 Sample structures, (a) LED with polarization dilution,(b) LED without polarization dilution (basic structure).
ß 2015 WILEYVCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

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a

922

(b)

1021

Ionizedacceptors

1019

-3

10

20

Posive polarizaoncharges

1018
1017

Acve
regionEB layer
gradedAlGaN

Charge density [cm ]-3Charge density [cm ](a)T. Yasuda et al.: Polarization dilution in a Ga-polar UV-LED

EB layer
uniform
AlGaN

1021
10

1019

Ionizedacceptors

1018
1017

P-layerPosive polarizaoncharges20

Acve
regionEB layer
uniform
AlGaN

P-layer

Figure 5 Distribution of positive polarization charges and ionized acceptors around EB layer, (a) LED with polarization dilution, (b)LED without polarization dilution

positive and negative charges, respectively. The reason forforming the energy spike (dip) simply indicates the result of
Coulomb’s force between the positive (negative) sheet charges
and the holes in the valence band. What we must pay attention tohere is the heights of the energy spike and the dip, 160 meV and
140 meV, respectively, similar to that of the heterobarrier at theAl0.2Ga0.8N/Al0.4Ga0.6N interface. This calculation resultclearly indicates that the inﬂuence of the polarization charges
on the energy band is large enough to affect device
performances. We carefully consider the energy spike due tothe polarization charges in the device design. For instance, suchan energy spike should be useful for carrier conﬁnement but
harmful for carrier transport in the same analogy as theheterobarrier. Thus, it is worthwhile to control and reduce thepolarization charges for better carrier transport. However, it is
difﬁcult to only reduce the polarization charges while keeping
the same mole fraction values at the interfaces.
3 Polarization dilution We then proposed a usefultechnique to “dilute” the polarization charges, as shown in
Fig. 3. Again, it seems impossible to just reduce thepolarization charges, but the charges could be neutralized byother charges with the opposite sign. In the case of theenergy spike in the p-type layer, negative charges arenecessary because the energy spike is caused by the positive
polarization charges. Unfortunately, the electron cannot be a
candidate in this case as the layer is p-type. Thus, anothernegative charge that can exist in a p-type layer is necessary,and that is an ionized acceptor. If the ionized acceptor
concentration is equal to or higher than the positive
polarization charge density, the positive polarization
charges will be completely neutralized. Note that a1  1013 cm2 sheet concentration is equivalent to a2  1020 cm3 volume concentration, assuming that the
sheet charges are distributed in one unit-cell thickness(0.5 nm). Such a high concentration is achievable as it isalready used in p-contact layers and tunnel junctions.
However, the high concentration could degrade the adjacentactive region in this case. Thus, we borrow the same conceptfrom polarization doping but in order to dilute thepolarization charges so that the required ionized acceptor
concentration will be reduced as much as possible. Here, theinterface with the positive polarization charges is compositionally graded with some thickness. In other words, theinterface thickness is expanded from 0 to the graded
thickness. For instance, if the interface with 1  1013 cm2

Figure 6 Comparison of band diagram (injected current density: 1 kA/cm2), (a) LED with polarization dilution, (b) LED without
polarization dilution.
ß 2015 WILEYVCH Verlag GmbH & Co. KGaA, Weinheim

www.pss-a.com

Original
Paper
Phys. Status Solidi A 212, No. 5 (2015)

923

Figure 7 Calculated LED characteristics. (a) J–V characteristic, (b) carrier injection efﬁciency.

is graded with 20 nm, the volume concentration is diluted
from 2  1020 cm3 to 5  1018 cm3, which is a required
value as a sufﬁcient ionized acceptor concentration and iseasily achieved by a conventional growth technique. Thus
the polarization dilution should be a very practical techniqueto suppress the inﬂuence of the positive polarization
charges. The differences between the polarization doping
and the polarization dilution are (1) the sign of thepolarization charges and (2) the requirement of impurity
doping. In the polarization doping case, the graded layer isformed at the interface generating the charges with the
opposite sign to the interested carriers, and the impurity
doping is not indispensable. On the contrary, in thepolarization dilution case, the graded layer is formed at
the interface generating the polarization charges with thesame sign as the interested carriers, and the impurity dopingis necessary to neutralize the charges.

(a)

4 Device simulation and experimental results
Finally, we theoretically and experimentally investigated
the impact of the polarization dilution in Ga-polar 350 nm
LEDs. Sample structures with and without (conventional)
the polarization dilution are shown in Fig. 4a and b,
respectively. A structural difference between the two LED
samples was (a) with and (b) without a compositional
graded p-AlGaN layer (Al composition from 0.2 to 0.4) atthe interface of the Al0.2Ga0.8N barrier/p-Al0.4Ga0.6N
electron blocking layer. Other than that, the layer structure
consists of a 2 mm AlN, a 1.5 mm n-Al0.2Ga0.8N, 3 nm GaN/
10 nm Al0.2Ga0.8N ﬁve quantum wells, the 10 nm u-AlGaN
last barrier layer, the 20 nm p-Al0.4Ga0.6N electron blocking
layer, a 100 nm p-Al0.2Ga0.8N, and a 30 nm p-GaN. Figures 5
and 6 show calculated proﬁles of the positive polarization
charge density and the ionized acceptor concentration, andcorresponding band-energy proﬁles, respectively. In the

(b)
8
7

Voltage [V]6with polarizaon diluon

5
4
3
2
1
0

Light Output Power[a.u.]

3.5

without
polarizaon diluon

3.0
2.5

without
polarizaon diluon

2.0
1.5
1.0
0.5

with polarizaon diluon
0

1000

2000

3000

4000

Current density[A/cm2]

5000

0.0

0

1000

2000

3000

4000

5000

Current density [A/cm2]

Figure 8 Measured LED characteristics. (a) J–V characteristic, (b) J–L characteristic.
www.pss-a.com

ß 2015 WILEYVCH Verlag GmbH & Co. KGaA, Weinheim

solidistatusphysica

pss

a

924

T. Yasuda et al.: Polarization dilution in a Ga-polar UV-LED

conventional LED case, the large polarization still remained
unneutralized at the active region/electron blocking layerinterface, as shown in Fig. 5, resulting in a large energy
spike, as shown in Fig. 6. On the other hand, in the LED withthe polarization dilution case, the positive polarization
charges are fully neutralized with the ionized acceptor andno energy spike is observed in the valence band at theinterface. In addition, while an ~100 meV reduction is
obtained at the valence band, no reduction of the potential
barrier is obtained at the conduction band. These results
should dramatically improve hole injection into the active
region. Figure 7 shows calculated J–V characteristics andinjection efﬁciency characteristics of the two LEDs. Due tothe reduction of the energy spike, the operating voltage isdecreased and the injection efﬁciency is increased in the
LED with the polarization dilution. These results suggestthat such a polarization-charge management is important toobtain superior performance in nitride-based UV-LEDs.
We then fabricated actual LEDs with the layer structures
shown in Fig. 4a and b. Figure 8 shows measured J–L–V
characteristics of the two LEDs. We obtained J–V
characteristics consistent with the calculated result. In themeantime, the LED with the polarization dilution showed asmaller light output power at low current injection region,
which is inconsistent with the calculated result. On thecontrary, in the high current injection region the light output
power of the LED with the polarization dilution was equal toor even higher than that of the LED without the polarization
dilution. At this time, the reason for the low light output
power is unclear. However, the result at high current
injection indicates the potential of the polarization dilution
and encourages us to investigate the reason in the future.

lower output power was simultaneously observed for somereason. However the result at high current injection
indicates the potential of the polarization dilution These
results suggested such polarization-charge management isimportant in the nitride semiconductors.

5 Conclusions In summary, the inﬂuence of thepolarization charges was calculated without the inﬂuenceof the band offset. The calculation result indicated thepolarization charges resulted in an energy spike of more than
100 meV, affecting the carrier transport. Then, we proposed
the polarization dilution to suppress the energy spike forbetter hole transport by using a Mg-doped graded AlGaN
layers in the UV-LEDs. The calculated results suggested
that the polarization dilution led to the lower operating
voltage and the higher injection efﬁciency. Our actual
350 nm LED also showed a lower operating voltage, whileß 2015 WILEYVCH Verlag GmbH & Co. KGaA, Weinheim

Acknowledgements This study was supported by MEXTGrants-in-Aid for Scientiﬁc Research Specially Promoted
Research (#25000011), Grant-in-Aid for Scientiﬁc Research (B)(#26286045), Nippon Sheet Glass Foundation for Materials
Science and Engineering, and The Asahi Glass Foundation.

References[1] D. Morita, M. Yamamoto, K. Akaishi, K. Matoba,
K. Yasutomo, Y. Kasai, M. Sano, S. Nagahama, and T.
Mukai, Jpn. J. Appl. Phys. 43, 5945 (2004).
[2] M. Shatalov, W. Sun, A. Lunev, X. Hu, A. Dobrinsky,
Y. Bilenko, J. Yang, M. Shur, R. Gaska, C. Moe, G. Garrett,and M. Wraback, APEX 5, 082101 (2011).
[3] H. Hirayama, S. Fujikawa, N. Noguchi, J. Norimatsu,
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206, 1176 (2009).
[4] W. Gotz, N. M. Johnson, J. Walker, D. P. Bour, and R. A.
Street, Appl. Phys. Lett. 68, 667 (1996).
[5] Y. Taniyasu, M. Kasu, and T. Makimoto, Nature 441, 325(2006).
[6] H. Amano, K. Nagamatsu, K. Takeda, T. Mori, H. Tsuzuki,
M. Iwaya, S. Kamiyama, and I. Akasaki, Proc. SPIE 7216,
72161B (2009).
[7] F. Bernardini, V. Fiorentini, and D. Vanderbilt, Phys. Rev. B
56, R10024 (1997).
[8] T. Takeuchi, S. Sota, M. Katsuragawa, M. Komori,H. Takeuchi, H. Amano, and I. Akasaki, Jpn. J. Appl. Phys.
36, L382 (1997).
[9] M. A. Khan, J. M. van Hove, J. N. Kuznia, and D. T. Olsen,
Appl. Phys. Lett. 58, 2408 (1991).
[10] D. Jena, S. Heikman, D. Green, D. Buttari, R. Cofﬁe,
H. Xing, S. Keller, S. DenBaars, J. Speck, U. K. Mishra, and
I. Smorchkova, Appl. Phys. Lett. 81, 4395 (2002).
[11] J. Simon, V. Protasenko, C. Lian, H. Xing, and D. Jena,
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[12] J. Verma, J. Simon, V. Protasenko, T. Kosel, H. G. Xing, and
D. Jena, Appl. Phys. Lett. 99, 171104 (2011).
[13] T. Yasuda, K. Yagi, T. Suzuki, T. Nakashima, M. Watanabe,
T. Takeuchi, S. Kamiyama, M. Iwaya, and I. Akasaki, Jpn.
J. Appl. Phys. 52, 08JJ05 (2013).

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Electron and hole accumulations at GaN/AlInN/GaN interfaces and conductive n-type
AlInN/GaN distributed Bragg reflectors

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REGULAR PAPER

Japanese Journal of Applied Physics 55, 05FD10 (2016)
http://doi.org/10.7567/JJAP.55.05FD10

Electron and hole accumulations at GaN/AlInN/GaN interfaces and conductive n-type
AlInN/GaN distributed Bragg reﬂectors
Shotaro Yoshida1*, Kazuki Ikeyama1, Toshiki Yasuda1, Takashi Furuta1, Tetsuya Takeuchi1,
Motoaki Iwaya1, Satoshi Kamiyama1, and Isamu Akasaki1,2
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan

2

Received December 31, 2015; accepted February 9, 2016; published online April 21, 2016
We investigated electron and hole accumulations at GaN/AlInN/GaN interfaces by Hall effect measurement. The InN mole fraction andtemperature dependences on the sheet carrier densities at the interfaces reveal that electrons and holes were induced by large positive andnegative polarization charges to satisfy the charge neutrality conditions, respectively. On the basis of the above results, we then designed anddemonstrated a low-resistity 10-pair Si-doped n-type AlInN/GaN distributed Bragg reﬂector (DBR) by using high Si doped and graded layers at the
GaN/AlInN interfaces. The low-resistity n-type AlInN/GaN DBR will reduce the resistance and the internal loss in blue vertical-cavity surface
emitting lasers. © 2016 The Japan Society of Applied Physics

1.

Introduction

AlInN is one of the important nitride-based alloys, which canbe lattice-matched to GaN, so that a high-quality AlInN=GaN
heterostructure can be achieved instead of a conventional
AlGaN=GaN. In optoelectronic devices, a thick AlInN
cladding layer1,2) and an AlInN=GaN distributed Bragg
reﬂector (DBR)3–5) were reported. One of the disadvantages
in growing thick AlInN had been the very low growth rate
(less than 0.2 µm=h6–8)) like GaInN alloys,9,10) but recently, a
relatively high growth rate, ∼0.5 µm, of the AlInN has beenobtained by using a combination of high growth temperatureand high In=III ratio.11) Now, a low-resistity current injectionin the vertical direction of the AlInN=GaN interfaces isrequired for various optoelectronic device applications. In
this case, a large hetero barrier (∼0.9 eV) at the AlInN=GaN
interfere in the lattice matching12) could lead to a highresistance, and a graded layer between AlInN and GaN
should be a solution as already proved in the AlGaAs=GaAs
system.13,14)
On the other hand, it is well known that large
polarization charges are induced at nitride-based heterointerfaces,15,16) and the lattice-matched AlInN=GaN is not
an exception.17) That is why a large electron accumulation
to positive polarization charges at the AlInN=GaN interface
is leveraged in a heterostructure ﬁeld-eﬀect transistor
structure.18,19) Then, a hole accumulation to negative polarization charges was observed at a GaN=AlGaN interface20)
in turn, and the AlN mole fraction dependence on the sheet
hole density was systematically investigated.21) Thus, thehole accumulation and its mole fraction dependence should
be observed at the AlInN=GaN interface in the samemanner, but such systematic investigations have not been
performed yet. The important thing is that not only thelarge hetero barrier height but also such large negative
polarization charges (and hole accumulations) at the AlInN=
GaN surely aﬀect the current injection across the AlInN=
GaN interface.
In this study, we ﬁrst investigated not only electron but
also hole accumulations to the GaN=AlInN=GaN interface
charges. The InN mole fraction and temperature dependenceson the sheet carrier densities at the interface were systematically investigated. On the basis of the results, we then

(a)

(b)
Fig. 1. (Color online) Sample structures (a) for the electron accumulation
at the bottom AlInN=GaN interface, and (b) for the hole accumulation at the
top GaN=AlInN interface.

designed and demonstrated conductive n-type AlInN=GaN
DBRs with high Si doped and graded layers.
2.

Experiment methods

All the samples were grown on c-plane sapphire substrates bymetalorganic vapor phase epitaxy. In this study, we prepared
two series of samples. One is a series of test structures forHall eﬀect measurement to investigate the electron and holeaccumulations at AlInN=GaN in the lattice matching and
GaN=AlInN interfaces with various InN mole fractions,
respectively. The other is a series of Si-doped 10-pair AlInN=
GaN DBRs to investigate vertical current injection throughthe DBRs.
Figures 1(a) and 1(b) show sample structures for theelectron accumulation at the interface of the AlInN=GaN (thebottom interface of the AlInN) and for the hole accumulation
at the interface of the GaN=AlInN interface (the top interface

05FD10-1© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD10 (2016)

S. Yoshida et al.

of the AlInN), respectively. The electron accumulation structure, the so-called two-dimensional electron gas (2DEG),
consists of a 20 nm low-temperature-deposited GaN buﬀer
layer (LT-GaN), a 3 µm undoped GaN, and a 90 nm undoped
AlInN. Six samples with various InN mole fractions from
0.13 to 0.23 in the AlInN were prepared. The hole accumulation structure, the so-called two-dimensional hole gas(2DHG), consists of the LT-GaN, the undoped GaN, a
60 nm undoped AlInN, a 10 nm undoped GaN and, a 40 nm
p-GaN contact layer. Four samples with various InN mole
fractions from 0.16 to 0.21 in the AlInN were prepared. X-ray
diﬀraction (XRD) patterns of 2θ=ω scan around the 0002
peak were measured to determine the InN mole fractions in
the AlInN layers. The sheet carrier densities at the interfaces
were measured by Hall eﬀect measurement. Electrodes forthe 2DEG and 2DHG samples were Cr=Ni=Au and Ni=Au,
respectively. In addition, we determined the temperaturedependence of sheet carrier density by low-temperature Halleﬀect measurement.
Regarding the 10-pair Si-doped AlInN=GaN DBRs, four
diﬀerent samples were prepared. After growing a 2 µm Sidoped GaN layer on the sapphire substrate, 10 pairs of about
45 nm AlInN and about 40 nm GaN were grown. The InN
mole fraction of the AlInN was adjusted to be 0.18 to belattice-matched to the GaN. Graded layers and high Si dopingat the top interfaces were investigated to see if better verticalcurrent injections through the DBRs were obtained. A
standard DBR was uniformly doped with Si (3 × 1018 cm−3
in GaN and 1 × 1019 cm−3 in AlInN), as shown in Fig. 2(a).
Note that the Si concentrations were measured by secondary
ion mass spectroscopy with reference samples of GaN:Si and
AlN:Si for calibration. A graded DBR was formed by using
9 nm graded AlGaInN layers at all the top interfaces of
the AlInN in the standard DBR, as shown in Fig. 2(b). The
compositional graded AlGaInN layer was grown by changing
Al, In, and Ga source ﬂow rates from a condition of AlInN
to a condition of GaN. A modulation-doped DBR consistedof 10 pairs of a 35.5 nm standard Si-doped (3 × 1018 cm−3)
GaN, a 4.5 nm highly Si-doped (1 × 1019 cm−3) GaN, a
4.5 nm highly Si-doped (5 × 1019 cm−3) AlInN, and a
35.5 nm standard Si-doped (1 × 1019 cm−3) AlInN, as shownin Fig. 2(c). Finally, a graded=modulation-doped DBR wasformed by a combination of the graded DBR and themodulation-doped DBR, as shown in Fig. 2(d). To measure
the vertical conductance of the 10-pair DBRs, a 60 µm mesa
was formed by etching the entire 10-pair DBRs to expose anunderlying n-GaN and depositing the top and bottom n-type
electrodes. Then, the current–voltage (I–V) characteristics ofthe mesa DBRs were measured at room temperature.
3.

Results and discussion

Figures 3(a) and 3(b) show the XRD patterns of (0002) 2θ=ωscan from the six AlInN=GaN structures for the investigationof electron accumulations [as shown in Fig. 1(a)] and thefour GaN=AlInN=GaN structures for the investigation of holeaccumulations [as shown in Fig. 1(b)], respectively. All thesamples show not only strong single peaks but also clear
Pendellösung fringes, meaning that the crystal and interface
qualities of AlInN=GaN layers in the lattice matching were
high. The InN mole fractions of the AlInN layers in the
structures were then estimated as indicated in the ﬁgures.

(a)

(b)

(c)

(d)
Fig. 2. (Color online) One pair of layer structures in the Si-doped AlInN=
GaN DBRs. (a) Standard DBR uniformly doped with Si (3 × 1018 cm−3 in
GaN and 1 × 1019 cm−3 in AlInN), (b) graded DBR formed with a 9 nm
graded AlGaInN layer at the top interface of the AlInN, (c) modulationdoped DBR consisting of a 35.5 nm standard Si-doped GaN, a 4.5 nm highly
Si-doped GaN, a 4.5 nm highly Si-doped AlInN, and a 35.5 nm standard Sidoped AlInN, and (d) graded=modulation-doped DBR formed by a
combination of the graded DBR and the modulation-doped DBR.

Next, we investigated the electrical properties of thesamples by Hall eﬀect measurement. Clear n-type conductivity was obtained in the samples shown in Fig. 1(a) even
without any Si doping. The measured sheet electron concentrations at the bottom interfaces as a function of InN mole

05FD10-2© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD10 (2016)

S. Yoshida et al.

(a)

-2

sheet electron concentration [cm ]

(a)

(b)
Fig. 3. (Color online) XRD patterns of (0002) 2θ=ω scan (a) from the six
AlInN=GaN structures [as shown in Figs. 1(a) and 1(b)] from the four GaN=AlInN structures [as shown in Fig. 1(b)]. The 34.5° peak is GaN, and theother peak is AlInN.

fraction are plotted with the red squares in Fig. 4(a). In
addition, theoretical positive polarization charge concentrations are also plotted in the same ﬁgure. The theoreticalpolarization charge density σ was calculated using thefollowing equation:22)
 ¼ Psp þ Ppz

c13
as  ae¼ Psp þ 2 e31 
e33 
;
ð1Þ
c33
ae
where Psp and Ppz are the spontaneous and piezoelectric
polarization charge densities, e31 and e33 are the piezoelectric
constants, c13 and c33 are the elastic constants, and ae andas are the lattice constants of the AlInN layers and theunderlying GaN layer, respectively. In this calculation, thematerial parameters listed in Table I were used, and all the
AlInN parameters were estimated assuming Vegard’s law.
We found that the measured sheet electron concentrations as
a function of InN mole fractions are in good agreement withthe theoretical polarization charge concentrations, meaningthat the electrons were accumulated to the positive polarization charges to satisfy the charge neutrality condition.
Also, no temperature dependence on the sheet electron concentration was clearly observed, as shown in Fig. 4(b),indicating that electrons were not directly generated fromnearby donor impurities.
Accordingly, p-type conductivity was clearly observed in
the samples shown in Fig. 1(b) even with Mg doping just forthe contact layer. The measured sheet hole concentrations at
the top interfaces of AlInN as a function of InN mole fraction were plotted with the blue triangles in Fig. 5(a).

10

14

10

13

10

12

InN
N : 0.13
InN
N : 0.17
InN
N : 0.21
4

6

8

10

12

14

-1

1000/T[K ](b)Fig. 4. (Color online) (a) Room-temperature sheet electron concentrations
as a function of InN mole fractions at the bottom interface of the AlInN=GaN
interface. The theoretical positive polarization charge densities are alsoplotted. (b) Temperature dependences on sheet electron concentrations at thebottom interface of the AlInN=GaN interface.

Table I. Material parameters used in the calculations.16,23)
GaN

AlN

InN

AlInN

Lattice constant (Å)
a

3.189

3.112

3.548

3.112x + 3.548(1 − x)

c

5.185

4.982

5.76

4.982x + 5.760(1 − x)−0.029−0.081
Psp
Piezoelectric constant (C=m2)−0.032−0.081x − 0.032(1 − x)

e31

−0.49

−0.6

−0.57

−0.6x − 0.57(1 − x)

e33

0.73

1.46

0.97

1.46x + 0.97(1 − x)

Spontaneous polarization (C=m2)

Elastic constant (GPa)
c13

100

127

108

127x + 108(1 − x)

c33

392

382

399

382x + 399(1 − x)

Similarly, theoretical negative polarization charge concentrations were plotted in the same ﬁgure. The measured sheet
hole concentrations are in good agreement with the theoretical polarization charge concentrations. Also, almost no
temperature dependence was observed in the sheet holeconcentrations as shown in Fig. 5(b). Thus, we concluded
that not just electrons but also holes are accumulated in thepolarization charges, resulting in high carrier concentrations
without standard impurity doping.
On the basis of the above results, a band diagram of
undoped GaN=AlInN=GaN is depicted in Fig. 6. This ﬁgure

05FD10-3© 2016 The Japan Society of Applied Physics

-2

sheet hole concentration[cm ]Jpn. J. Appl. Phys. 55, 05FD10 (2016)
5x10

13

4x10

13

3x10

13

2x10

13

1x10

13

S. Yoshida et al.

calculation
measurement

0
0.12

0.16

0.20

0.24

InN mole fraction-2

sheet hole concentration[cm ]

(a)
10

14

Fig. 7. (Color online) Conduction band proﬁles of one pair in thestandard DBR, the graded DBR, the modulation-doped DBR, and thegraded=modulation-doped DBR.

10

13

InN : 0.17
InN : 0.20
10

12

4

6

8

10

12

14

1000/T(b)
Fig. 5. (Color online) (a) Room-temperature sheet hole concentrations as
a function of InN mole fractions at the top interface of the GaN=AlInN
interface. The theoretical negative polarization charge densities are alsoplotted. (b) Temperature dependences on sheet hole concentrations at the top
interface of the GaN=AlInN interface.
Fig. 8. (Color online) I–V characteristics of the standard DBR, the graded
DBR, the modulation-doped DBR, and the graded=modulation-doped DBR.

Fig. 6. (Color online) Band diagram of undoped GaN=AlInN=GaN.

clearly suggests that almost no Si doping is required at
the bottom interfaces of AlInN while heavy Si doping is
indispensable at the top interfaces of AlInN to obtain
conductive n-type DBRs. This is the reason why we tried
to utilize not only graded layers but also heavy Si doping in
the AlInN=GaN DBRs. Actually, the large negative polarization charges could be neutralized with the ionized Si
atoms,24) so the amount of Si doping should be close to thatof the polarization charge. For instance, the polarization sheet
charge density at the Al0.83In0.17N=GaN is calculated andmeasured to be about 3 × 1013 cm−2, which corresponds to3 × 1019 cm−3 volume concentrations in a 10 nm layer. Thus,
the 10 nm area around the GaN=AlInN interface should be
doped with Si up to such a high concentration.
We then calculated the band proﬁles of the four AlInN=
GaN DBRs described above by using a commerciallyavailable device simulator. Figure 7 shows the conduction

band proﬁles of one pair in the four AlInN=GaN DBRs.
Note that the Fermi level was set to be zero in Fig. 7. The
calculated results clearly show that the graded layer and the
Si modulation-doped at the top GaN=AlInN interfaces arevery useful for reducing the potential height around the top
interfaces.
Figure 8 shows the I–V characteristics of the four 10-pairSi-doped AlInN=GaN DBRs. The modulation-doped DBR
and the graded=modulation-doped DBR showed clear ohmic
characteristics with very low resistances. The graded=
modulation-doped DBR showed the lowest resistance among
the samples as the calculated results suggested. The values
resistance is 16 Ω, but includes not only a vertical componentof the DBR but also a lateral component of the underlying
GaN, so the vertical resistance of the DBR must be smaller
than 16 Ω. On the other hand, almost no current passed
through the standard DBR within 3 V of the applied voltage
range. Interestingly, the graded DBR showed some current
ﬂowing but a high resistance. The above results suggest thatthe layer with high Si doping is more eﬀective than thegraded layer to obtain the conductive n-type AlInN=GaN
DBRs, meaning that the impact of the polarization charges on
the electrical properties is huge in nitride-based materials. So
far, laser operations of nitride-based vertical cavity surface
emitting lasers (VCSELs) with undoped AlInN=GaN DBRs

05FD10-4© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FD10 (2016)

S. Yoshida et al.

were reported,25–27) and the undoped DBRs were notconductive at all, so that intracavity contacts must be used,
resulting in high resistance and internal loss. As a result, sucha conductive nitride-based DBR will lead to high-performance blue VCSELs, like the already commercialized infrared
VCSELs with the conductive DBRs.28)
4.

Conclusions

We demonstrated the low-resistity 10-pair Si-doped AlInN=
GaN DBR by using the graded layer and high Si doping at
the top interfaces of the AlInN layers. The reason for the highSi doping is to neutralize the large negative polarization
charges at the interface which resulted in the large potential
barrier. We also demonstrated electron and hole accumulations induced by the polarization charges at the interfaces
by investigating InN mole fractions and temperature dependences on the sheet carrier densities measured by Halleﬀect measurement.
Acknowledgments

This study was supported by the MEXT Program for theStrategic Research Foundation at Private Universities,
2012–2016, Grants-in-Aid for Scientiﬁc Research Specially
Promoted Research (No. 25000011), and Grant-in-Aid for
Scientiﬁc Research (B) (No. 26286045).

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High-photosensitivity AlGaN-based UV heterostructure-field-effect-transistor-type
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REGULAR PAPER

Japanese Journal of Applied Physics 55, 05FJ04 (2016)
http://doi.org/10.7567/JJAP.55.05FJ04

High-photosensitivity AlGaN-based UV heterostructure-ﬁeld-effect-transistor-type
photosensors
Akira Yoshikawa1,2*, Yuma Yamamoto1, Takuya Murase1, Motoaki Iwaya1, Tetsuya Takeuchi1,
Satoshi Kamiyama1, and Isamu Akasaki1,3
1Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Asahi Kasei Corporation, Fuji, Shizuoka 416-8501, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan2*E-mail: yoshikawa.ac@om.asahi-kasei.co.jp
Received November 5, 2015; accepted January 4, 2016; published online April 5, 2016
We report the development of high-performance AlGaN/AlN heterostructure-ﬁeld-effect-transistor-type (HFET) photosensors with a p-type GaN
optical gate and detection wavelengths that are restricted to 220–280 nm. These photosensors employ a two-dimensional electron gas induced
at the hetero-interface between Al0.6Ga0.4N and Al0.5Ga0.5N as a highly conductive channel. In addition, a p-type GaN optical gate is employed
to deplete a channel. Consequently, we obtained a high photosensitivity of over 4 ' 103 A/W and an externally low dark current density of
approximately 5 ' 10%10 A/mm at a source–drain voltage of 3 V. We also determined that the detection range of light wavelength in these HFET
photosensors can be controlled by controlling the AlN molar fraction in the AlGaN channel layer. The results are very promising for thedevelopment of completely solar-blind high-performance photosensors with high photosensitivity.
© 2016 The Japan Society of Applied Physics

1.

Introduction

Group-III nitride semiconductors are widely used as lightemitting diodes (LEDs)1–4) and laser diodes.5,6) Since theband gap of AlGaN ternary alloys covers a broad range from
3.39 to 6.2 eV at room temperature,7,8) these alloys are
suitable as UV emitters and in photosensor applications.9–13)
Recently, high-performance UV LEDs with an external
quantum eﬃciency close to 10% have been developed forsterilization and medical applications.14–18) To make these
UV LEDs practical, photosensors are indispensable.
The photosensors are required to be small, durable, lowcost, solar-blind and highly sensitive, and possess a highrejection ratio. Furthermore, the most important feature is thatthe detection wavelength must be freely controlled. The
detection of speciﬁc wavelengths leads to detection of ﬂames
or organic matter, which is advantageous. In the case of aﬂame sensor, the detection of a speciﬁc wavelength is very
important because almost all wavelengths from a ﬂame areburied under sunlight and only a 250 to 280 nm peak isseparated from sunlight.19)
Until now, photomultiplier tubes (PMTs) or Si-based
photodiodes with a visible light cut-oﬀ ﬁlter have been used
as UV sensors.20,21) However, even with its high photosensitivity, there are issues regarding PMT miniaturization,such as cost, heat resistance, and high operating voltage.
Furthermore, the Si-based photodiodes is up to four orders ofmagnitude lower than that of PMTs. Nitride semiconductors
have the potential to solve the above-mentioned problem of
PMTs and Si photodiodes. Thus far, we have reported highsensitivity AlGaN=GaN heterostructure-ﬁeld-eﬀect-transistor-type (HFET) photosensors with a p-type GaN opticalgate.22,23) These photosensors employ a two-dimensional
electron gas (2DEG) induced at the hetero-interface between
AlGaN and GaN as a highly conductive channel. Althoughthese sensors exhibited a high photosensitivity of 105 A=W,
they detected all wavelengths lower than 365 nm.
We assumed that the detection wavelength is in accordancewith the band gap of the channel layer. Therefore, we tried tocontrol the detection wavelength by changing the AlN molar
fraction of the AlGaN channel layer.
In this study, we employed an Al0.5Ga0.5N channel layer
because a band gap of 4.5 eV was required to detect 280 nm
UV light. However, the fabrication of an AlGaN=AlGaN
HFET is a big challenge. One reason for this is the prereaction between Al and NH3.24) The higher the Al molar
fraction of AlGaN, the more drastic the reaction. Therefore,
investigating optimum growth condition is important. The
other reason is that many cracks are generated because of thetensile stress between AlGaN and the GaN substrate.25) An
AlN substrate or an AlN template substrate is eﬃcient at
solving this problem because it has a smaller lattice size than
GaN. Moreover, there are other problems; only a 1-in. AlN
substrate could be obtained and the crystallinity of an AlN
template substrate is worse than that of a GaN template
substrate.26,27)
For the above reasons, until Akita et al. reported an
Al0.51Ga0.49N-based HEMT, there were only a few reportsof AlGaN=AlGaN-HEMT and the AlN molar fraction x of
AlxGa1−xN was lower than 0.4.28)
In this study, we fabricated high-performance AlGaN=
AlGaN HFET photosensors with a p-GaN optical gate on
high-quality AlN template substrates for the detection of
wavelengths that are restricted to 220–280 nm. We will also
show that the AlN molar fraction of the AlGaN channel layer
is the key to controlling the detection wavelength.
2.

Experimental Procedure

Figure 1 shows a schematic view of an AlGaN HFET-type
photosensor with a p-type GaN optical gate. The device wasgrown using metalorganic vapor phase epitaxy on a sapphire(0001) substrate. The growth condition of the AlN template
was based on a previous report.29) After the growth of a 1.8µm AlN layer, a 320-nm unintentionally doped Al0.5Ga0.5N
layer, a 35-nm unintentionally doped Al0.6Ga0.4N barrier
layer, and a 100-nm Mg-doped p-GaN layer with a Mg
concentration of 2 × 1019 cm−3 was grown. The Al molar
fraction x of AlxGa1−xN was determined using 2θ=ω scans and

05FJ04-1© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FJ04 (2016)

A. Yoshikawa et al.

8 μm

(a)

(a)

1 x 10-5λ=250 nm

2 μm

Photocurrent, Ip (A/mm)

u-Al0.6Ga0.4N-35 nm
Source2DEG

Ti/Al/Ti/Au

u-Al0.5Ga0.5N-320 nm

Photocurrent, Ip

1 x 10-6

p-GaN100 nm

Drain
Ti/Al/Ti/Auu-AlN-1.8 μm

1 x 10-7

Under room lighting
Under 1sun (AM 1.5G)

1 x 10-8

Dark current, Id

1 x 10-9(0002) Sapphire substrate
1 x 10-10(b)

1 x 10-5(b)

Photocurrent, Ip

8 μm

Mesa

Photocurrent, Ip (A/mm)

1 x 10-62 μm

100 μmDark current, Id

1 x 10-7

1 x 10-8

1 x 10-9
SourceDrain
1 x 10-10
0.0

30 μm

1.0

2.0

3.0

Source-Drain voltage, VSD (V)Fig. 1. (Color online) Schematic views of the device structure: (a) crosssectional view and (b) plan view.

reciprocal space mapping measurements with X-ray diﬀraction. The Mg acceptors were activated by annealing at 850 °C
for 5 min in N2. Mesa isolation was performed using Cl2
inductively coupled plasma (ICP) etching. The p-type GaN
optical gate was etched using ICP etching, except in the lightdetection area. Then, Ti (30 nm)=Al (100 nm)=Ti (20 nm)=Au (150 nm) was deposited as the source and drain electrodeson the u-Al0.6Ga0.4N barrier layer and annealed at 810 °C in
N2. The length and width of the p-GaN optical gate were 2 and100 µm, respectively, and the interval between the source anddrain electrodes was 8 µm. No anti-reﬂection coating was used
for any of the devices in this study.
The photocurrent at each light wavelength (220–1100 nm)
was measured using a semiconductor parameter analyzer(HP-4155B) and a spectroscope with a Xe lamp (HAL-320)
at room temperature (RT). Furthermore, the photocurrent wasmeasured under room lighting conditions and air mass (AM)
1.5G pseudo sunlight (HAL-320), respectively. The photosensitivity was calculated from the photocurrent, lightdetection area, and irradiation density of the spectroscope,
respectively. In this study, we employed a monochromatic
light irradiation density of approximately 45 µW=cm2 and alight-detection area of 200 µm2. For the evaluation of thesensor as a ﬂame sensor, a Bunsen burner was used to
produce a ﬂame and was placed 10 cm from the sensor.

Fig. 2. (Color online) ISD–VSD characteristics of the Al0.6Ga0.4N=
Al0.5Ga0.5N HFET photosensors at RT in the VSD range from 0 to 3 V withand without irradiation from a UV lamp (λ = 250 nm, irradiation density =45 µW=cm2). (a) With and (b) without an optical p-type GaN gate.

3.

Results and Discussion

Figure 2 shows the ISD–VSD characteristics of the HFET
photosensors (a) with and (b) without an optical p-type GaN
gate. The photocurrent from UV light irradiation at 250 nm(irradiation density: 45 µW=cm2) and the dark current at RT
were measured. An extremely low dark current of 5 × 10−10
A=mm was obtained using the HFET in the presence of an
optical p-GaN gate when VSD was 3 V, as shown in Fig. 2(a).
This dark current was almost the same as the measurement
limit of our measurement system. In contrast, a large darkcurrent of approximately 1 × 10−6 A=mm was observed whenusing HFET photosensors without an optical p-GaN gatewhen VSD was 3 V, as shown in Fig. 2(b).
In addition, we compared the photocurrents of photosensors with an optical p-GaN gate under room lighting
conditions and an AM 1.5 G solar simulator (HAL-320) at
a light intensity of 1 sun. Consequently, we measured aphotocurrent of 5.7 × 10−10 A=mm under room lighting
conditions and 1.1 × 10−9 A=mm under 1 sun (AM 1.5 G).
These values are in an order of magnitude equal to that ofthe dark current. The results indicate that the photosensors

05FJ04-2© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FJ04 (2016)

A. Yoshikawa et al.

1 x 105

Table I. Photosensitivity of a Si UV sensor, a photoelectric tube, a Si
photodiode, a GaN UV sensor, a photomultiplier tube, and the photosensor
used in this study.

Al0.2Ga0.8N/GaN [22]1 x 105

1 x 103

S(A/W)

Photoresponsivity, S(A/W)

1 x 104

1 x 1031 x 1011 x 1021 x 10-1
1 x 101200

1 x 100

300 400
λ(nm)Peak wavelength(nm)

Photosensitivity(A=W)

This work250

3 × 103

Si UV sensor

254

3 × 10−2

Photoelectric tube

254

2 × 10−2

GaN photodiode254

6 × 10−2

Photomultiplier tube

400

2 × 105500
1 x 10-5

In dark room

Al0.6Ga0.4N/Al0.5Ga0.5N/AlN

Photocurrent with a flame1 x 10-1
Under room lighting
1 x 10-2
0

200

400

600

800

1000

1200

Wavelength, λ(nm)Fig. 3. (Color online) Relationship between the photosensitivity with a
VSD of 3 V and the wavelength of irradiating light. The inset is focused
between 200 and 500 nm. The circular plots show the photosensitivity of
Al0.6Ga0.4N=Al0.5Ga0.5N-HFET with a VSD of 3 V. The triangular plots refer
to Al0.6Ga0.4N=Al0.5Ga0.5N-HFET with a VSD of 3 V under room lighting
conditions. The square plots refer to Al0.2Ga0.8N=GaN-HFET with a VSD
of 5 V.

Photo current, Ip (A/mm)

1 x 10-6

1 x 10-7

1 x 10-8

Photocurrent without a flame1 x 10-9

1 x 10-10
0.0

1.0

2.0

3.0

Source-Drain voltage, VSD (V)were almost non-responsive to sunlight. However, they werecapable of detecting very weak light with a wavelength of
250 nm. Therefore, this device is a completely solar-blind
UV sensor.
Figure 3 shows the monochromatic photosensitivity of
HFET-type photosensors from 200 to 1100 nm. The circular
plots show the photosensitivity of Al0.6Ga0.4N=Al0.5Ga0.5NHFET (with an operating voltage of 3 V between the source
and the drain) obtained using a Xe lamp and a spectrometer.
The photosensors can achieve a high external photosensitivity exceeding 103 A=W. The inset in Fig. 3 is focused
between 200 and 500 nm. The absorption edge and peak
photosensitivity wavelengths of the photosensors wereapproximately 280 and 250 nm, respectively. The rejection
ratio of the photosensors was greater than 104.
The triangular plots show the photosensitivity under room
lighting conditions. Almost identical absorption edge andpeak photosensitivity was obtained. In contrast, for the
darkroom measurement, little absorption was observed
between 280 and 450 nm, compared with that between 520
and 1100 nm for room lighting measurements. This absorption may be caused by the p-GaN layer.
The square plots show the photosensitivity of Al0.2Ga0.8N=
GaN-HFET obtained at VSD = 5 V.22) The absorption edge
and peak photosensitivity wavelengths of the photosensors
were approximately 410 and 300 nm, respectively. The
photocurrent and dark current are higher compared withthose of Al0.6Ga0.4N=Al0.5Ga0.5N-HFET. From simulation
results (SiLENSe), we determined the sheet carrier densityof Al0.6Ga0.4N=Al0.5Ga0.5N-HFET to be 1 × 106 cm−2, while
it was 3 × 1012 cm−2 for Al0.2Ga0.8N=GaN. Therefore, weFig. 4. (Color online) ISD–VSD characteristics of the Al0.6Ga0.4N=
Al0.5Ga0.5N-HFET photosensors at RT under a VSD that varied from 0 to 3 Vwith ﬂame irradiation from a Bunsen burner.

speculate that these results were caused by the diﬀerence inthe carrier concentration of 2DEG.
From the above, these results indicate that the absorption
edge and peak photosensitivity wavelengths of the photosensors were controlled by the AlN molar fraction of the
AlGaN channel layer. In addition, the photocurrent and darkcurrent is probably adjustable by altering the AlN molar
fraction and the thickness of the AlGaN barrier layer.
We also compared the photosensitivity of the photosensors
with those of commercially available photosensors. Table I
shows the monochromatic photosensitivity of each photosensor. From this table, the photosensors used in this study
possessing photosensitivity that is superior to those of a Si
UV sensor (KYOSEMI KPDU400F-2), a photoelectric tube
(Hamamatsu Photonics R765), a Si photodiode (Hamamatsu
Photonics S1087), and a GaN photodiode (KYOSEMIKPDU27HQ1). Moreover, the results imply that the photosensors used in this study have a photosensitivity that isequivalent to that of photomultiplier tubes.
We also investigated the performance of the photosensors
as ﬂame sensors. The photocurrent was measured in a
darkroom, under room lighting conditions, and at 1 sun (AM1.5G) using a parameter analyzer both with a ﬂame andwithout a ﬂame. The result in a dark room is shown in Fig. 4.
Table II summarizes the results for the photocurrent with andwithout a ﬂame, respectively. For all conditions, a sensitivity

05FJ04-3© 2016 The Japan Society of Applied Physics

Jpn. J. Appl. Phys. 55, 05FJ04 (2016)
Table II.
a ﬂame.

A. Yoshikawa et al.

Photocurrent of HFET when VSD was 3 V both with and without
Photocurrent (A=mm)

Darkroom
Under room lighting
Under 1 sun (AM 1.5G)

with a ﬂame

without a ﬂame

1.12 × 10−7

9.98 × 10−10−7

5.68 × 10−10−71.06 × 10−9

1.34 × 10
1.09 × 10

to noise (S=N) of over 102 was obtained. These results
indicate that the photosensors can be used as a ﬂame sensor
under practical environmental conditions.
4.

Summary and conclusion

We fabricated high-performance Al0.6Ga0.4N=Al0.5Ga0.5NHFET-type photosensors on AlN for the detection of 220
to 280 nm wavelengths. It was shown that the AlN molar
fraction of the AlGaN channel layer is the key to controllingthe detection wavelength. The measured photosensitivity of
3.4 × 103 A=W greatly surpasses those of commerciallyavailable photosensors and is comparable with those ofphotomultiplier tubes. These photosensors are completely
solar-blind and exhibit high performance as a ﬂame sensor
under both room lighting conditions and 1 sun (AM 1.5G).
Acknowledgments

This study was partially supported by the Program for theStrategic Research Foundation at Private Universities,
2012–2016, supported by the MEXT, a MEXT Grant-inAid for Specially Promoted Research No. 25000011, andthe MEXT Grant-in-Aid for Scientiﬁc Research A No.
15H02019. The authors would like to thank Ph.D. T.
Morishita, T. Okumura, and S. Ushida for discussions andadvice.

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Correlation between Device Performance and Defects in GaInN-Based Solar Cells

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Applied Physics Express 5 (2012) 082301
http://dx.doi.org/10.1143/APEX.5.082301

Correlation between Device Performance and Defects in GaInN-Based Solar Cells
Mikiko Mori1 , Shinichiro Kondo1 , Shota Yamamoto1 , Tatsuro Nakao1 , Takahiro Fujii1 , Motoaki Iwaya1 ,Tetsuya Takeuchi1 , Satoshi Kamiyama1 , Isamu Akasaki1;3 , and Hiroshi Amano2;3
1

Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan3
Akasaki Research Center, Nagoya University, Nagoya 464-8603, Japan
2

Received June 20, 2012; accepted June 26, 2012; published online July 13, 2012
We investigated the correlation between the device performance and defects, such as V-shaped pits and threading dislocations, in GaInN-basedsolar cells. To realize high-performance GaInN-based solar cells with a high open-circuit voltage and fill factor, it is essential to realize a low pit
density of less than 107 cm 2 . In this study, we were unable to observe clear evidence of any effect of the threading dislocation density in the GaN
underlying layer. # 2012 The Japan Society of Applied Physics

itride semiconductor material systems including
AlN, GaN, InN, and their alloys are widely used
as light-emitting diodes (LEDs) and laser diodes.
Since the bandgap of GaInN ternary alloys covers a broad
range from 0.65 to 3.43 eV, these alloys are suitable for solar
cell applications.1) Using a combination of GaInN alloys,
one may theoretically design a multijunction photovoltaic
device with a bandgap spanning the infrared to ultraviolet
regions of the solar spectrum. This makes it possible toachieve high-eﬃciency multijunction solar cells.2)
So far, we have succeeded in fabricating GaInN-basedsolar cells.3–7) By improving the crystal quality of GaInN
using freestanding c-plane GaN substrates3) and applying
GaInN superlattice (SL) structures,4) the conversion eﬃciency has been improved.
In general, GaN ﬁlms on sapphire substrates covered
with a low-temperature buﬀer layer8,9) have a high densityof threading dislocations ranging from 108 to 1010 cm2 . In
contrast, GaN substrates grown by hydride vapor phase
epitaxy10) with a low dislocation density of 105 to 107 cm2
are commercially available. In addition, GaInN ﬁlms exceeding the critical thickness on a GaN template have beenreported to induce the generation of misﬁt dislocations andthe growth of V-shaped pits.11) There have been many
reports discussing the correlation between device performance and crystal defects in GaInN-based LEDs, laser
diodes, and so forth. For instance, because dislocations act asnonradiative recombination centers, reducing the dislocation
density in GaN is very important for the realization of higheﬃciency nitride-based LEDs.12,13) In addition, it has beenreported that dislocations have a strong impact on the device
lifetime of violet laser diodes.14) The impact of V-shaped
pits and misﬁt dislocations on the performances of LEDs andlaser diodes have also been reported.15,16)
In contrast, we reported that the leakage current in GaInNbased solar cells increases with increasing density of Vshaped pits in the devices.3) However, there have been few
reports discussing the correlation between device performance and the density of defects such as dislocations andV-shaped pits in GaInN-based solar cells. Understanding thefactors limiting the performance of such devices is essential
for realizing high-performance GaInN-based solar cells.
In this study, we investigated the correlation between
the device performance and defects in GaInN-based solar
cells.

N

E-mail address: 113434038@ccalumni.meijo-u.ac.jp

(a)

(b)

(c)

Fig. 1. Schematic view of the sample structures of nitride-based solar
cells: (a) type I, (b) type II, and (c) type III.

All samples were grown by metalorganic vapor phase
epitaxy. c-Plane sapphire and freestanding c-plane GaN wereused as the substrates. In the case of the sapphire substrate,
GaN underlying layers were grown using a low-temperaturebuﬀer layer. The threading dislocation densities of GaN on
sapphire and freestanding GaN substrates are approximately
3  108 and less than 107 cm2 , respectively. Trimethylindium, trimethylaluminum, trimethylgallium, triethylgallium,
and ammonia were used as the source gases. Figures 1(a)–1(c)show schematic views of the sample structures. We prepared
three types of samples on freestanding GaN and sapphire
substrates. In the type I samples [Fig. 1(a)], a single GaInN
active layer was used as an active layer. In the type II samples[Fig. 1(b)], GaInN-based SLs were used as an active layer. In
the type III samples [Fig. 1(c)], to suppress the formation ofpits and the threading of misﬁt dislocations,17) we used thick
GaInN-based SLs as the active region instead of GaInN SLs(underlying SLs). Table I summarizes the detailed structureof the GaInN active layer and substrate of each sample. Pits
with a diameter of more than 50 nm were counted when
determining the pit density using plan-view scanning electron
microscope images. The thickness and InN molar fraction in
the GaInN active layer were determined by X-ray diﬀraction
 diﬀractions. We also con2=! scans of (0002) and (2024)
ﬁrmed the thickness of each layer in some samples by crosssectional transmission electron microscopy. The density ofV-shaped pits ranged from 1:1  105 to 8:8  108 cm2 asshown in Table I. The threading dislocation density also
ranged over two orders of magnitude from 107 to 109 cm2 .
Semitransparent Ni (5 nm)/Au (5 nm) ohmic contacts
to p-GaN and Ti/Al/Ti/Au ohmic contacts to n-GaN wereformed by electron beam evaporation. No antireﬂection
coating was used for all the devices. The solar cell characteristics were analyzed using a solar simulator (air mass
1.5G, irradiation strength 155 mW/cm2 ). All measurements
were carried out at room temperature (RT).

082301-1# 2012 The Japan Society of Applied Physics

Appl. Phys. Express 5 (2012) 082301
Table I.

M. Mori et al.

Device structure parameters of GaInN-based solar cells on sapphire and GaN substrates. In this table, the pit density of each sample is also given.

Sample No.

Substrate

Underlying SL

Active layer

Pit density (cm2 )

1

GaN

None

Ga0:88 In0:12 N single layer(200 nm)

8:0  108

2

GaN

None

50 pairs Ga0:83 In0:17 N (3 nm)/Ga0:93 In0:07 N (0.6 nm) SLs

8:7  108

3

GaN

None

50 pairs Ga0:83 In0:17 N (3 nm)/Ga0:93 In0:07 N (3 nm) SLs

1:5  107

4GaN

10 pairs Ga0:90 In0:10 N: Si (3 nm)/GaN: Si (3 nm) SLs

50 pairs Ga0:83 In0:17 N (3 nm)/Ga0:93 In0:07 N (0.6 nm) SLs

1:0  106

5

Sapphire

10 pairs Ga0:92 In0:08 N: Si (1.8 nm)/GaN: Si (2.2 nm) SLs

30 pairs Ga0:90 In0:10 N (3 nm)/Ga0:90 In0:10 N (0.6 nm) SLs

3:2  105

6

Sapphire

10 pairs Ga0:90 In0:10 N: Si (2.2 nm)/Ga0:95 In0:05 N: Si (2.2 nm) SLs+
10 pairs Ga0:94 In0:06 N: Si (1.8 nm)/GaN: Si (2.2 nm) SLs

50 pairs Ga0:82 In0:18 N (3 nm)/Ga0:88 In0:12 N (0.6 nm) SLs

1:7  108

7

Sapphire

10 pairs Ga0:94 In0:06 N: Si (2.2 nm)/GaN: Si (1.8 nm) SLs

30 pairs Ga0:88 In0:12 N (3 nm)/Ga0:93 In0:07 N (0.6 nm) SLs

1:1  105

2.5

70

1.5
1.0

3

60

2

50
40
30

0.5
0.0

GaN sub.
Sapphire sub.

20
5

10

6

7

8

10
10
10
-2
Pit density [cm ]910

JSC [mA/cm ]FF [%]

2.0
VOC [V]

480

GaN sub.
Sapphire sub.

10

5

10

6

1
GaN sub.
Sapphire sub.

7

8

10
10
10
-2
Pit density [cm ]

0

9

10

5

10

(b)

(a)

2

6

7

8

10
10
10
-2
Pit density [cm ]

9

10

(c)

Fig. 2. (a) VOC of GaInN-based solar cells as a function of pit density. (b) FF as a function of pit density. (c) JSC as a function of pit density.

Figures 2(a)–2(c) show the open-circuit voltage (VOC ), ﬁll
factor (FF), and short-circuit current density (JSC ) of the
GaInN-based solar cells on freestanding GaN and sapphire
substrates as a function of the pit density at the surface.
From these ﬁgures, JSC was found to be independent of thepit density. In contrast, there is a signiﬁcant dependenceof VOC and FF on the pit density. No signiﬁcant diﬀerence
was found between JSC , VOC , and FF for the samples grown
on GaN and those grown on sapphire substrates, meaningthat these characteristics are not aﬀected by the threading
dislocation density in the underlying layer.
Figures 3(a)–3(c) summarize the leakage current density(Jleak ) at a reverse bias voltage of 3 V, the shunt resistance(RSH ), and the series resistance (RS ), respectively, of eachsample as a function of pit density. We calculated RS and
RSH from the slope of the I–V curves.18) From these results,
Jleak and RSH were found to be strongly dependent on the pit
density. In Fig. 3(a), the results for Jleak are ﬁtted to theformulapﬃﬃﬃﬃﬃﬃ
ð1Þ
Jleak ¼ 4  108 Dp ;where Dp and 4  108 A/cm are the pit density and ﬁtting
parameter, respectively, as shown by the blue dotted line
in Fig. 3(a). There have been several reports that suggestthat the origin of the leakage current in an LED structure isrelated to V-shaped pits, although the carrier conduction
mechanism is still unclear at present.19,20) In this experiment,

Jleak exhibits sublinear dependence on the pit density and isalmost proportional to the square root of Dp . Because RSH
should be expressed as the reverse bias voltage divided
by Jleak , in Fig. 3(b) we inserted the theoretical line (bluedotted line) given by
RSH ¼

VR
VR
pﬃﬃﬃﬃﬃﬃ :¼
Jleak
4  108 Dp

ð2Þ

In this case, VR is the reverse bias voltage of 3 V. In thisﬁgure, the ﬁtting curve closely matches the experimental
results. From Figs. 2(a), 2(b), and 3(b), RSH should exceed
104  cm2 and Dp should be less than 107 cm2 to obtainsuitable values of VOC and FF.
Next, we compared these results with the solar cell characteristics of other material systems. Table II summarizes
RSH , VOC , and the bandgap of each material in refs. 21–23.
According to this table, to realize a high VOC in a GaInNbased solar cell, RSH should be one to two orders of magnitude higher than that in the other material systems.
Note that VOC is deﬁned by the formula24)Iph  ðVOC =RSH Þ
nkB T
ln
VOC ¼
þ1 ;
ð3Þ
q
I0
where n, kB , T , Iph , and I0 are the nonideality factor,
Boltzmann constant, absolute temperature, photocurrent, andreverse saturation current, respectively. From this formula,
VOC is signiﬁcantly dependent on RSH . In contrast, there was no082301-2# 2012 The Japan Society of Applied Physics

Appl. Phys. Express 5 (2012) 082301

M. Mori et al.

-2

6

10

10

4

10

GaN sub.
Sapphire sub.

GaN sub.
Sapphire sub.

-3

10

2

2

-4

10

3

10

5

RS [Ω cm ]10

RSH [Ω cm ]2

Jleak [A/cm ]GaN sub.
Sapphire sub.

4

10

-5

10

2

10

1

10
3

10

-6

10

5

10

6

7

8

10
10
10
-2
Pit density [cm ]

9

0

5

6

10

10

7

8

10
10
10
-2
Pit density [cm ]

9

10

5

10

6

7

8

10
10
10
-2
Pit density [cm ]

9

10

(c)

(b)

(a)

10

Fig. 3. (a) Jleak of GaInN-based solar cells at a reverse bias voltage of 3 V as a function of pit density. (b) RSH as a function of pit density. (c) RS as afunction of pit density.

Table II. RSH , VOC , and bandgap of each material in this study and inrefs. 21–23.

Material system

Bandgap energy
(eV)

RSH
( cm2 )1.12

6:7  102

0.56

GaInP

1.96

2

1:1  10

1.33

CdS/CdTe23)

1.4

1  1030.84

2.8

7  105

1.95

2.8

1  104

1.8

2.8

2:15  103

0.93

GaInN

87 (2005) 595.
3) Y. Kuwahara, T. Fujii, Y. Fujiyama, T. Sugiyama, M. Iwaya, T. Takeuchi,

VOC
(V)

22)

Si21)

2) H. Hamzaoui, A. S. Bouazzi, and B. Rezig: Sol. Energy Mater. Sol. Cells

4)

5)

6)

7)

correlation between RS and pit density as shown in Fig. 3(c). In
general, considering the equivalent circuit model of a solar cellunder illumination, the resistive component is expressed as22) qðV þ RS IÞ
V þ RS I
I ¼ Iph  I0 exp
;
ð4Þ
1 
nkB T
RSH
where I0 is the reverse saturation current. Equation (4) showsthat JSC was strongly aﬀected by the value of RS . Therefore,
JSC is independent of the pit density.
In conclusion, we investigated the correlation between thedevice performance and defects in GaInN-based solar cells.
To realize high-performance GaInN-based solar cells witha high VOC and FF, it is essential to realize a low Dp andhigh RSH , which should be lower than 107 cm2 and exceed
104  cm2 , respectively. In contrast, no signiﬁcant correlation could be conﬁrmed between these device performances
and the threading dislocation density in the GaN underlying
layer. Therefore, reduction of the pit density is essential forthe realization of high-performance GaInN-based solar cells.
Acknowledgments This study was partially supported by the New Energy
and Industrial Technology Development Organization (NEDO) Project,‘‘Research and Development on Innovative Solar Cells’’ and the Program
for the Strategic Research Foundation at Private Universities, 2012–2016,
supported by the Ministry of Education, Culture, Sports, Science andTechnology.

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